Rare-earth ion doped chalcogenide
waveguide amplifiers
Kunlun Yan
May 2018
A thesis submitted for the degree of Doctor of Philosophy of
The Australian National University
Laser Physics Centre
Research School of Physics and Engineering
College of Physical and Mathematical Sciences
The Australian National University
© Copyright by Kunlun Yan 2018
All Rights Reserved
Acknowledgements
I would like to express my sincere gratitude to my supervisors A/Prof. Steve Madden,
Dr. Rongping Wang and Prof. Barry Luther-Davies, for their continuous support of my
Ph.D study and related research. This thesis could not have been possible without their
guidance and support. For everything you’ve done for me, thank you.
My sincere thank also goes to Dr. Khu Vu who guided me and gave me suggestions
from the first day I entered Laser Physics Centre till now. I would like to thank Dr. Duk
for his help in waveguide fabrication, and thank Dr. Zhingyong Yang for glass
fabrication.
I thank our technical staff Sukanta for his help in film deposition, Craig and John for
their assistance in fixing equipments used in the lab. I also thank our department
administrator Sonia for her kind help. Also, I want to acknowledge my colleagues in
ANFF for providing the fabrication facility.
I am very grateful to have a group of friend, Xin Gai, Ivy, Joseph, Ting Wang, Niko,
Zhisong Qu, they make my student life so much more enjoyable.
Last, I would like to thank my family. Thank my parents for their tireless love and
support over these years. Thank my beloved wife Pan for her encouragement, constant
support and love. Thank you.
Publications
1. Kunlun Yan, Khu Vu, Rongping Wang, and Steve Madden, "Greater than 50%
inversion in Erbium doped Chalcogenide waveguides", Optics Express, 24, 20,
23304-23313, 2016.
2. Kunlun Yan, Khu Vu, and Steve Madden, “Internal gain in Er-doped As2S3
chalcogenide planar waveguides”, Optics Letters, 40, 5, 796-799, 2015.
3. Kunlun Yan, Khu Vu, Zhiyong Yang, Rongping Wang, Sukanta Debbarma,
Barry Luther-Davies, and Steve Madden. “Emission properties of erbium-doped
Ge-Ga-Se glasses, thin films and waveguides for laser amplifiers”, Optical
Materials Express, 4, 3, 464-475, 2014.
4. Kunlun Yan, Rongping Wang, Khu Vu, Robert Elliman, Kidane Belay and
Barry Luther-Davies. “Photoluminescence in Er-doped Ge-As-Se Chalcogenide
Thin Films”, Optical Materials Express, 2, 9, 1270-1277, 2012.
5. Rongping Wang, Kunlun Yan, Mingjie Zhang, Xiang Shen, Shixun Dai, Xinyu
Yang, Zhiyong Yang, Anping Yang, Bin Zhang, and Barry Luther-Davies,
“Chemical environment of rare earth ions in Ge28.125Ga6.25S65.625 glass-ceramics
doped with Dy3+”, Applied Physics Letters, 107, 161901, 2015.
6. Rongping Wang, Kunlun Yan, Zhiyong Yang, Barry Luther-Davies, “Structural
and physical properties of Ge11.5As24S64.5-xSe64.5-(1-X) glasses”, Journal of Non-
crystalline Solids, 427, 16-19, 2015.
7. Vu K, Yan K, Jin Z, Gai X, Choi D, Debbarma S, Luther-Davies B, Madden S.
“Hybrid waveguide from As2S3 and Er-doped TeO2 for lossless nonlinear
optics”, Optics Letters, 38, 11 1766-1768, 2013.
8. Xinyu Yang, Mingjie Zhang, Kunlun Yan, Liyuan Han, Qin Xu, Haitao Liu,
Rongping Wang, "Controllable Formation of the Crystalline Phases in Ge–Ga–S
Chalcogenide Glass-Ceramics", Journal of American Ceramic Society,
doi:10.1111/jace.14492
I
Abstract
Chalcogenide glass waveguide devices have received a great deal of attention
worldwide in the last few years on account of their excellent properties and potential
applications in mid-infrared (MIR) sensing and all-optical signal processing. Waveguide
propagation losses, however, currently limit the potential for low power nonlinear
optical processing, and the lack of suitable on chip integrated MIR sources is one of the
major barriers to integrated optics based MIR sensing. One approach to overcome the
losses is to employ rare-earth ion doped waveguides in which the optical gain can
compensate the loss, in such a way that the conversion efficiency of nonlinear effects is
increased significantly. For infrared applications, the long wavelengths potentially
attainable from rare-earth ion transitions in chalcogenide hosts are unique amongst glass
hosts. New rare-earth ion doped chalcogenide sources in the MIR range could benefit
molecular sensing, medical laser surgery, defence etc. Despite these promising
applications, until now, no one has succeeded in fabricating rare-earth ion doped
chalcogenide amplifiers or lasers in planar devices.
This work develops high quality erbium ion doped chalcogenide waveguides for
amplifier and laser applications. Erbium ion doped As2S3 films were fabricated using
co-thermal evaporation. Planar waveguides with 0.35 dB/cm propagation loss were
patterned using photolithography and plasma etching on an erbium ion doped As2S3
film with an optimised erbium ion concentration of 0.45x1020 ions/cm3. The first
demonstration of internal gain in an erbium ion doped As2S3 planar waveguide was
performed using these waveguides. With different film deposition approaches,
promising results on intrinsic lifetime of the Er3+ 4I13/2 state were achieved in both ErCl3
doped As2S3 films (2.6 ms) and radio frenquency sputtered Er3+:As2S3 films (2.1 ms),
however, no waveguide was fabricated on these films due to film quality issues and
photopumped water absorption issues.
The low rare-earth ion solubility of As2S3 is considered the main factor limiting its
performance as a host. Gallium-containing chalcogenide glasses are known to have
good rare-earth ion solubility. Therefore, a new glass host material, the Ge-Ga-Se
system, was investigated. Emission properties of the bulk glasses were studied as a
function of erbium ion doping. A region between approximately 0.5 and 0.8 at% of Er3+
ion was shown to provide sufficient doping, good photoluminescence and adequate
lifetime to envisage practical planar waveguide amplifier devices. Ridge waveguides
II
based on high quality erbium ion doped Ge-Ga-Se films were patterned. Significant
signal enhancement at 1540 nm was observed and 50 % erbium ion population
inversion was obtained, in waveguides with Er3+ concentration of 1.5×1020 ion/cm3. To
the Author’s knowledge, this is the highest level of inversion ever demonstrated for
erbium ions in a chalcogenide glass host and is an important step towards future devices
operating at 1550 nm and on the MIR transitions of erbium ions in chalcogenide glass
hosts. Photoinduced absorption loss caused by upconversion products in the waveguides
is the remaining hurdle to achieving net gain. Further research is needed to find suitable
compositions that possess high rare-earth ion solubility whilst avoiding the detrimental
photoinduced losses.
Table of contents
Abstract ................................................................................................................................... I
Abbreviations .......................................................................................................................... i
Chapter 1 Introduction ........................................................................................................... 1
1.1 General introduction ..................................................................................................... 1
1.1.1 Integrated optics ..................................................................................................... 1
1.1.2 Highly nonlinear waveguide device limitations .................................................... 2
1.1.3 MIR sensing opportunity ....................................................................................... 4
1.2 Chalcogenide glasses and their applications ................................................................ 6
1.2.1 Chalcogenide glasses ............................................................................................. 6
1.2.2 Example applications and limiting performance factors of chalcogenide
glasses ............................................................................................................................. 8
1.3 Rare-earth ion doped chalcogenide glasses ................................................................ 12
1.3.1 Introduction .......................................................................................................... 12
1.3.2 History of rare-earth ion based chalcogenide laser and amplifier devices .......... 13
1.3.3 Rare-earth ion doped chalcogenide materials ...................................................... 15
1.3.4 Potential issues for rare-earth ion doped chalcogenide glasses ........................... 19
1.4 Problem under study ................................................................................................... 21
1.5 Outline of this work .................................................................................................... 22
Chapter 2 Rare-earth ions in chalcogenide glasses .............................................................. 25
2.1 Introduction to rare-earth ions .................................................................................... 25
2.2 Properties and characterization of rare-earth ions in glasses...................................... 27
2.2.1 Lifetime of rare-earth ions in an excited energy level ......................................... 27
2.2.2 Absorption and emission cross-section ............................................................... 30
2.2.3 Judd-Ofelt theory ................................................................................................. 33
2.3 Energy migration processes in rare-earth ion doped materials .................................. 37
2.3.1 Ion-Ion interactions .............................................................................................. 37
2.3.2 Up-conversion process..........................................................................................38
2.3.3 Energy transfer......................................................................................................40
2.3.4 Cross relaxation ....................................................................................................41
2.4 Concentration quenching .............................................................................................42
2.5 Solubility of rare-earth ion in glasses ..........................................................................43
2.6 Conclusion ...................................................................................................................44
Chapter 3 Co-thermal evaporation of erbium ion doped As2S3 waveguides ........................45
3.1 Diarsenic trisulfide: As2S3 ...........................................................................................45
3.1.1 Background of As2S3 ............................................................................................45
3.1.2 Properties of As2S3 ...............................................................................................46
3.2 Erbium ion doped As2S3 ..............................................................................................48
3.2.1 Research in erbium ion doped As2S3 glasses ........................................................48
3.2.2 Fabrication of bulk As2S3 glass ............................................................................50
3.2.3 Erbium ion doped As2S3 film deposited using co-thermal evaporation ...............51
3.3 Characterization of erbium ion doped As2S3 films .....................................................55
3.3.1 Raman spectrum of erbium ion doped As2S3 film deposited using thermal
evaporation ....................................................................................................................55
3.3.2 Optical properties of erbium ion doped As2S3 film ..............................................56
3.4 Up-conversion properties of erbium ion doped As2S3 film .........................................61
3.5 High temperature thermal post-treatment on erbium ion doped As2S3 film ...............67
3.6 Concentration effects in erbium ion doped As2S3 films ..............................................69
3.7 Conclusion ...................................................................................................................71
Chapter 4 Erbium ion doped As2S3 waveguide amplifiers ...................................................73
4.1 Basic theory and simulation methods ..........................................................................73
4.1.1 Maxwell’s equations for optical waveguides .......................................................73
4.1.2 Simulation of optical waveguide ..........................................................................74
4.2 Waveguide design .......................................................................................................75
4.2.1 Waveguide geometry ........................................................................................... 75
4.2.2 Mode properties of the hybrid integrated waveguide structure ........................... 80
4.3 Simulation of amplification performance of erbium ion doped As2S3 waveguide
amplifier ........................................................................................................................... 82
4.3.1 Rate equation for an erbium ion doped waveguide system ................................. 82
4.3.2 Simulation of erbium ion doped As2S3 waveguide amplifiers ............................ 85
4.4 Waveguide fabrication ............................................................................................... 88
4.4.1 Waveguide fabrication introduction .................................................................... 88
4.4.2 Erbium ion doped As2S3 waveguide fabrication ................................................. 93
4.5 Characterization of erbium ion doped As2S3 waveguides .......................................... 95
4.5.1 Propagation loss and absorption of erbium ion doped As2S3 waveguides .......... 96
4.5.2 Amplification measurements of erbium ion doped As2S3 waveguides ............... 98
4.6 Co-thermal evaporation of neodymium ion doped As2S3 ........................................ 102
4.6.1 Neodymium ion doped As2S3 film deposited using co-thermal evaporation .... 103
4.6.2 Characterization of the neodymium ion doped As2S3 film ................................ 104
4.7 Conclusion ................................................................................................................ 110
Chapter 5 Different approaches to fabricate erbium ion doped As2S3 waveguides ........... 111
5.1 Co-thermal evaporation of ErXn (X=S/Cl) with As2S3 ............................................ 111
5.1.1 Co-thermal evaporation of ErCl3 with As2S3 ..................................................... 112
5.1.2 Co-thermal evaporation of Er2S3 with As2S3 ..................................................... 115
5.2 Radio-frequency (RF) sputtering of erbium ion doped As2S3 films ......................... 117
5.2.1 Introduction to RF sputtering chalcogenide glass films .................................... 117
5.2.2 Erbium ion doped As2S3 film deposited using RF sputtering ............................ 118
5.2.3 Characterisation of RF sputtered erbium ion doped As2S3 ................................ 119
5.2.4 Pump-induced PL intensity decay in RF sputtered erbium ion doped As2S3
films ............................................................................................................................ 122
5.3 Conclusion ................................................................................................................ 128
Chapter 6 Erbium ion doped Ge-Ga-Se waveguides ......................................................... 131
6.1 Erbium ion doped Ge-Ga-Se bulk glasses .................................................................131
6.1.1 Background on Ge-Ga-Se glasses ......................................................................131
6.1.2 Erbium ion doped Ge-Ga-Se glasses synthesis and characterisation .................133
6.1.3 Optical properties of erbium ion doped Ge-Ga-Se glasses .................................135
6.1.4 PL and PL lifetime of erbium ion doped Ge-Ga-Se bulk glasses .......................136
6.1.5 Absorption and emission cross-section of erbium ion doped Ge-Ga-Se bulk
glasses ..........................................................................................................................138
6.1.6 Erbium ion doped Ge25Ga10Se65 bulk glass summary ........................................142
6.2 Erbium ion doped Ge-Ga-Se film deposition and characterization ..........................143
6.2.1 Erbium ion doped Ge-Ga-Se film deposition .....................................................143
6.2.2 Erbium ion doped Ge-Ga-Se film characterisation ............................................144
6.3 Erbium ion doped Ge-Ga-Se film modification and waveguide characterisation.....148
6.3.1 Modification on erbium ion doped Ge-Ga-Se film deposition ...........................148
6.3.2 Erbium ion doped Ge-Ga-Se waveguides characterization ................................151
6.4 Optical enhancement and photoinduced absorption of erbium ion doped Ge-Ga-Se
waveguides ......................................................................................................................153
6.5 Conclusion .................................................................................................................158
Chapter 7 Conclusions and recommendations ....................................................................159
7.1 Conclusions ...............................................................................................................159
7.2 Recommendations for future work ............................................................................160
References ...........................................................................................................................163
i
Abbreviations
Abbreviations
AO: acousto-optic
ASE: amplified spontaneous emission
BPM: beam propagation method
CCD: charge-coupled device
CVD: chemical vapour deposition
CW: continuous wave
DI water: deionisation water
DPSD: differential power spectral density
EDS: energy dispersive x-ray spectroscopy
ESA: excited state absorption
ETU: Energy transfer up-conversion
FDM: finite difference method
FDTD: finite-difference time-domain method
FEM: The finite element method
FSO: Free-Space Optical Communication
FWHM: full width at half maximum
FWM: four wave mixing
GLS: gallium lanthanum sulfide
ICP: inductively coupled plasma
LAN: local area network
MAN: metropolitan area network
MIR: mid-infrared
MoL: method of lines
MPR: multi-phonon relaxation
NA: numerical aperture
NIR: near-infrared
NLO: nonlinear optic
OPA: optical parametric amplifier
OPO: optical parametric oscillator
OSA: optical spectrum analyser
PL: photoluminescence
PLE: photoluminescence excitation spectroscopy
PML: perfect matching layer
ii
QCL: quantum cascade lasers
RF: radio-frequency
RIE: reactive ion etching
RMS: Root-mean-square
RMSE: root-mean-square error
SC: supercontinuum source
SEM: scanning electron microscope
TE mode: transverse electric mode
Tg: glass transition temperature
TM: transition metal
TM mode: transverse magnetic mode
TOX: thermally oxidized silicon
TV: television
UV: ultraviolet
V/A: volts/amps
WAT: weak absorption tail
WDM: wavelength division multiplexing
1
Chapter 1
Introduction
1.1 General introduction
1.1.1 Integrated optics
Integrated optics has been a hot topic as electronics has been slowly supplemented by,
and, in some cases, even replaced by, optics particularly in telecommunications systems
since the 1980s. The original inspiration for this concept came from electronic
integrated circuit technology [1], which displayed phenomenal advances in
technological capability and cost reduction and is now truly ubiquitous. Thus, it is easy
to understand that the main objectives of integrated optics are essentially similar to
those of integrated electronics, namely that bulk components can be eliminated, and
multiple functionalities can be integrated on one stable robust, mass producible, low
cost device. Like optical fibres, the basic concept in optical integrated circuits is light
confinement. In planar waveguides, light propagates within a region embedded in a
planar substrate or within a portion of a film deposited upon a planar substrate due to
the higher refractive of channel to that of the substrate/surrounds. This thin film type
optical circuit aims to perform multiple functions by integrating laser light sources,
functional components such as switches, modulators, filters, amplifiers, interconnecting
waveguides, and detectors all on single substrate. By integration, mass producible, low
cost, compact, stable and highly functional optical systems can be achieved [2].
To attain the full suite of desired functions in an integrated optical device, the
functions and capabilities that have to be realised in a planar geometry include:
• On chip laser source (continuous wave (CW) and mode locked) capability
• Optical amplification
• Large and ultrafast second and/or third order optical nonlinearity
• Electronic control of refractive index for tuning/switching/modulation
• Acousto-optic effects for widely tunable filters
• Tight bending capability to make compact devices
• Low loss interfaces to standard optical fibres for the spectral range of interest
• Wide transmission range
2
• Stable photosensitivity for performance tuning and Bragg grating devices
• Light detection capability
Realising all of these different functionalities in a single material system is
extremely challenging (InP integrated technology perhaps coming closest [3]), and so
hybridisation methods are likely required.
It is now that clear one key challenge is finding a host material which can enable the
most critical subset of the functions mentioned above and that has a clear route to
hybridise with other materials easily to obtain the rest, and also that has compatibility
with current electronic fabrication technology. Additionally, it should also enable high
yield, high volume, low cost production. Currently, several different materials are used
in commercial high performance integrated devices, such as silicon dioxide, silicon
(oxy) nitride, polymers, lithium niobate, silicon, and indium phosphide. Each material
has its advantages and disadvantages. Despite the plethora of existing platforms and the
massive investments currently being made, there are several technology gaps still
hindering the applicability of integrated optics at a broad level i.e. beyond telecoms.
These include good nonlinear optic (NLO) performance, MIR transmission, having a
wide choice of gain wavelengths and usable acousto-optic (AO) capability.
Chalcogenide glasses, with high refractive index, high third order nonlinearity, MIR
transparency, low phonon energies, and high AO coefficients offer opportunity in all
these areas. Therefore, it is worthwhile considering chalcogenide glasses as a major
component in a new hybrid platform for integrated optics.
1.1.2 Highly nonlinear waveguide device limitations
Highly nonlinear waveguides fabricated from chalcogenide glasses are proving
extremely effective for all-optical signal processing of high-speed telecommunications
signals e.g. [4-9]. To achieve the most efficient processors, the natural response is to
maximise the nonlinear response of the waveguide which requires the most nonlinear
material along with a waveguide design that gives the smallest mode area. When
combined these lead to a large nonlinear coefficient, , of the waveguide, given by
equation (1.1):
𝛾 = 2𝜋𝑛2 𝜆𝐴𝑒𝑓𝑓⁄ (1.1)
where n2 is the third order optical nonlinearity of the material, denotes the wavelength
of operation and Aeff is the mode area of the waveguide [10]. Recently reported
3
Ge11.5As24Se64.5 (at%) nanowires ≈500×630 nm in cross-section achieved mode
effective areas of 0.25 m2 and values of 130-150 W-1m-1 [11], which is amongst the
highest nonlinearity reported for any materials system and is the highest for a glass
based system [4]. However, nanowires experience significant propagation losses
because of the strong interaction of the optical fields to the etched sidewalls even when
they have been fabricated using optimized processes that produce the lowest roughness
(<2 nm Root-mean-square (RMS)) [11]. Typically the losses achieved for highly
nonlinear chalcogenide nanowires are >1.5 dB/cm whilst significantly lower values
(<0.2 dB/cm) have been achieved in As40S60 (at%) [12, 13] with ~1 m2 mode effective
area and 0.42 dB/cm in a sub 1 micron squared Ge-Sb-S [14] rib waveguides. However,
as the nonlinear phase shift for a given amount of input power also depends on the
effective length (i.e. the loss limited interaction length) then often better performance is
obtained using longer lower loss devices as the scales almost linearly with Aeff (i.e.
width/height) whereas the loss scales essentially exponentially with waveguide
dimension [15, 16]. These losses strongly limit nonlinear performance. For example, a
commonly used nonlinear process in all-optical processing is four wave mixing (FWM).
Based on the analytical expression for FWM conversion efficiency from Batelgelj [17],
with the assumption of negligible pump depletion, and in the low power regime at the
optimum length, the converted internal signal power in a waveguide is given by
equation (1. 2):
𝑃𝐷𝐹𝑊𝑀 =4
27𝑃𝑠 [
𝛾𝑃𝑝𝑢𝑚𝑝
𝑎]2
(1. 2)
where PDFWM is the idler power generated by a signal power Ps and a degenerate pump
power Ppump, is the waveguide nonlinear parameter, is the pump-signal frequency
spacing dependent phase mismatch efficiency parameter, and a is the loss of waveguide.
Clearly, from this equation there is an inverse square dependence of the overall idler
conversion power on the propagation loss of waveguide, so a significant improvement
can be expected at the output with waveguide loss reduction [17]. One key means of
reducing loss is to compensate it with gain, and in fact performing FWM in an optical
amplifier is advantageous and results in large increases in efficiency [18-23].
Thus, a method of overcoming the waveguide losses is desirable to enhance the
efficiency of all-optical processing. One approach is to employ a rare-earth ion doped
nonlinear waveguide in which optical gain compensates the loss. The 1.55 μm emission
arising from the 4I13/2→4I15/2 transition in the trivalent state of the erbium ion has been
4
widely used to create optical fibre and waveguide amplifiers in oxide glass hosts [24-
29]. Hence it should, in principle, be possible to use gain to compensate the optical
losses in a chalcogenide waveguide as a route to improving the efficiency of all-optical
processing.
1.1.3 MIR sensing opportunity
One of the major challenges for MIR sensing is the availability of either high power
broadband sources, or widely tunable lower power laser devices. The best known route
to obtain high power and high brightness broad band sources is supercontinuum
generation [30, 31]. However, to generate broadband MIR supercontinuum, it is optimal
to use a pump around 4 microns with femtosecond pulses [32-34]. Currently no such
commercial mode locked laser exists and so there is an opportunity for a suitable device
there, especially one integrated with an SC generator. In terms of tunable devices,
currently, quantum cascade lasers (QCLs) and optical parametric oscillator/amplifier
(OPO and OPAs) serve as the main workhorse in MIR photonics. For QCLs, the huge
amount of heat produced in CW mode laser consumes up to 70 % of the injected
electrical power, with the given relatively small active area ~100 μm2, brings challenges
to realise higher power single-mode infrared light [35], although QCLs with output
power up to several watts are readily available [36]. However, the narrow tuning ranges
of QCLs and the fact they cannot be mode locked to generate high peak power fs pulses
set limitations to their applications. OPOs that offer wide tunability and multiwatt
output power are now commercially available, but the excitation required by parametric
generation needs to be narrow linewidth and linear polarized, which are significant
constraints on the pump [37]. Also, OPOs are always expensive and physically large.
Chalcogenide ceramic lasers based on Cr2+:ZnSe are commercially available within
a broad wavelength range in the near- and mid-infrared. But issues arise from thermal
lensing and quenching from multiphonon emission is becoming a challenge when
higher power output is required [35]. A bulk optics mode locked Cr2+:ZnSe laser with
pulse duration from 40-150 fs, providing over 1 W output power in the range of 2.1-2.6
μm is available from IPG Photonics, and lasers with longer operation wavelength up to
5 μm are available based on liquid nitrogen cooled Fe2+:ZnSe [38].
Rare-earth ion doped devices as an alternative to the technologies discussed above
have important characteristics. For example, the absorption and emission wavelengths
of rare-earth ions are relatively insensitive to the host material and correspond to readily
5
available diode laser wavelengths for low cost pumping. The metastable state lifetimes
of rare-earth ions are long, and the quantum efficiency can be high. Besides this,
different rare-earth ions offer the opportunity to get a wide range of emission
wavelengths and often wide tuning ranges which almost cover the entire spectrum from
visible to 12 μm [35, 39, 40]. Rare-earth ions also offer an excellent path for mode
locking to produce high peak power ultra-short pulses on account of their large emission
bandwidth, and mode locking in rare-earth ion doped fibre lasers has been extensively
demonstrated and today dominates the commercial femtosecond laser market in the
guise of the Yb3+ doped silica fibre laser.
These excellent properties lead to a significant opportunity for rare-earth ions in
MIR laser/amplifier applications where there are few other options. In addition, there
are numerous telecommunications based opportunities for rare-earth ion doped devices,
not yet fully exploited by the now ubiquitous silica fibre based erbium ion doped optical
amplifier. As an example, some of the most popular rare-earth ions with their potential
lasing wavelengths in the infrared are listed in Table 1. 1.
Table 1.1 Rare-earth ions with their potential lasing wavelength from [41] with updated
details in italics.
Rare-earth ions IR emission wavelengths (µm)
Nd3+ 0.786, 0.919, 1.08, 1.37 [42]
Er3+ 0.822, 0.86, 0.987 [43, 44], 1.54 [45], 2.7, 3.5, 4.5 [46],
Tm3+ 1.21, 1.45, 1.81, 2.35 [47], (3.8, 5.38 [48])
Ho3+ 0.76, 0.91 [44], 1.2, 2.9, 3.9 [49], (2.1 [50], 3 [51])
Pr3+ 1.3, 1.6, 2.9, 3.4, 4.5, 4.8, 4.9, 7.2 [52, 53], (5.5, 7.6 [54])
Dy3+ 1.3, 1.8, 2.3, 4.3 [55], (3.0, 3.2, 4.5, 5.5, 7.6 [54])
Tb3+ 3.0, 4.8, 8.0 [46], (10.5 [54])
Yb3+ 0.975, 1.02, 1.14 [56]
It is well known that silica optical fibre becomes opaque quickly beyond 2 μm, and
although laser operation up to 3.9 μm has been demonstrated in fluoride glass hosts,
beyond 3 μm the multiphonon relaxation rates of fluoride glasses compete with the
sharp-line luminescence. Laser emission centred at 3.5 μm with the longest wavelength
emission at 3.78 μm was recently realised in an erbium ion doped ZBLAN fibre laser
6
[57], and 3.9 μm fibre laser was realised in a Ho3+-doped ZBLAN fibre pumped by a
dye laser at 640 nm [58], but no laser with longer operation wavelength than this was
reported on ZBLAN so far. With rather lower phonon energies, chalcogenide glasses,
therefore, offer the prospect of attaining longer wavelength higher efficiency lasing.
As noted above, there are clearly some important opportunities that could be
addressed using rare-earth ion doped chalcogenide glasses, especially when the proven
planar integration capabilities of chalcogenide are factored in. Therefore, rare-earth ion
doped chalcogenide planar waveguides were chosen as the topic of this thesis.
1.2 Chalcogenide glasses and their applications
1.2.1 Chalcogenide glasses
A chalcogenide is a chemical compound that contains one or more of the chalcogen
elements from Group 16 of the Periodic Table (normally referred as sulfur, selenium,
and tellurium but excluding oxygen) covalently bonded to network formers such as
arsenic, antimony, germanium, silicon, etc. The heavy elements are relatively weakly
covalently bonded and this provides chalcogenide glasses with some important optical
and physical properties. For example, low bond energy leads to [4, 59]:
• Chalcogenide glasses’ band gaps located in the visible or near infrared.
• Relatively low glass transition temperature (Tg), from ~100-400 °C.
• Extended optical transparency (12 µm in sulfide to 20 µm in telluride).
• High polarizability leading to large second and third order optical nonlinearity.
• Low hardness offering high acousto-optic coefficients.
Typical transmission spectra of several classes of glass including chalcogenide
glasses are shown in Figure 1.1. Due to large atomic masses and relatively weak bond
strengths resulting in low phonon energies, chalcogenide glasses have a long
wavelength cut-off that lies in the MIR [60]. The transparency edge is 12 μm for sulfide
based glasses, 15 μm for selenide glasses and 20 μm for telluride glasses, whilst silica
becomes opaque at 4 μm and the transparency of ZBLAN drops quickly beyond 6 μm.
This offers enormous opportunities to chalcogenide glasses in MIR applications.
7
Figure 1.1 Transmission spectra for several glasses (optical path length of about 2-3 mm)
[61].
There are a number of optical properties chalcogenide glasses possess that have
made them of interest for several decades. High linear refractive index (2-3.5) and low
absorption suggest that compact optical circuits made from sub-wavelength, single
mode waveguides can be fabricated for telecommunications and MIR science. The low
phonon energies present in chalcogenide glasses permit radiative transitions for rare-
earth ion dopants that are not possible in silica and phosphate glasses due to
multiphonon quenching, making them promising rare-earth ion host materials for IR
amplifiers and lasers. The high photosensitivity present in many chalcogenide glasses
enables direct optical patterning of planar waveguides, gratings and other photonic
devices. The high optical non-linearity displayed by many chalcogenide glasses
combined with low multi-photon absorption has been used for some of the most
successful all-optical signal processing demonstrations of telecommunications signals
[4, 9, 11, 62-66]. Finally, the wide range of glass-forming and compositions allow
material properties to be fine-tuned for specific applications. In Table 1.2, some basic
properties of chalcogenide, tellurite, silica and fluoride glasses are shown for
comparison.
8
Table 1.2 Basic properties of chalcogenide glass in context with other optical glasses [67].
Property Chalcogenide Tellurite Silica Fluoride
Refractive index (RI) 2.1-3.5 1.8-2.3 1.46 1.5
Thermal coefficient of RI /10-6 °C-1 ~20-60 [68]
(@4.515 µm)
~-20 [69]
(@1.5 µm)
~12 [70]
(@1 µm)
~-10 [71]
(@1.53 µm)
Highest phonon energy / cm-1 ≥250 ~800 ~1000 ~500
Transmission range / µm 0.8-20 0.4-5 0.2-2.5 0.2-7.0
Bandgap / eV 1-3 ~3 ~10 ~10
Nonlinear refractive index (n2) / m2
W-1
3-10×10-18 3-6×10-19 1-3×10-20 ~10-21
Glass transition temperature (Tg) /
°C
100-400 300-450 ~1000 ~300
Thermal expansion /10-7 °C-1 ~140-250 120-170 ~5 ~150
Density / kg m-3 4500 5500 2200 5000
Note: the thermal property values for chalcogenide glass change rapidly with glass
composition, the values shown in this table are for representative Ge-As-Se glasses.
Amongst the properties chalcogenide glasses possess, low phonon energies, long
wavelength transparency and relatively high rare-earth ion solubility make chalcogenide
glasses attractive in exploring laser sources in the near and MIR range. Based on the
definition in ISO 20473:2007E (Optics and photonics — Spectral bands), the MIR
refers to the spectral region from 3 μm up to 50 μm [72]. However, the wavelength
range of 2–25 μm covers the important atmospheric windows, and the molecular
fingerprints of numerous gases, liquids and solids [39], and has gained widespread
acceptance as a working definition that is used hereafter in this thesis. Laser sources
working in this range have great potential applications in various fields, and these
applications will be illustrated in the following section.
1.2.2 Example applications and limiting performance factors of chalcogenide
glasses
Besides the applications in integrated optics, nonlinear waveguide devices and MIR
sensing mentioned in the general introduction part, there are other potential applications
under intensive study as well.
9
(a) Remote molecular sensing
The MIR between 400 and 4000 cm−1 (25–2.5 μm) covers the important range for
vibrational spectroscopy, making MIR technology attractive for chemical or
biochemical sensors. The applications here span a huge range of fields through
medicine, agriculture, climate and atmospheric monitoring, industrial process control,
border security, petrochemicals, water quality monitoring, etc. MIR spectroscopy has
been known and used for decades in some of these areas, but field deployment has been
impossible due to the lack of low power, rugged, low cost instruments. This could be
overcome with planar integration.
One of the most challenging but essential parts for integrated infrared optical
sensing is the light source. For bulk optics based devices the thermal Globar© source
can offer light with a huge spectral range, but the brightness is not sufficient enough for
applications where high spatial resolution or long distance propagation is required [73],
and it is not easily planar integratable. Quantum cascade lasers (QCL) provide
diffraction-limited MIR beams with high brightness, but the tuning range is quite
narrow meaning multiple sources need to be flip chip bonded increasing cost and
complexity.
Wideband amplified spontaneous emission (ASE) sources (potentially with multiple
dopants and cascade emission) are another option for white light sources. For example,
the spectral range from 1540 nm to >2340 nm at −20 dB was achieved in Tm3+–Ho3+
co-doped alumina-rich silica core fibre [74]. An ASE emission peak at 1880 nm with a
3-dB spectral width of almost 2000 nm was generated with Bi3+–Tm3+ co-doped
lithium-alumino-germano-silicate (LAGS) core fibre [75]. Praseodymium ion in
chalcogenide fibre has also been demonstrated as an ASE source, offering an emission
bandwidth spanning 3 μm to 5 μm at -20 dB [40, 76] with the potential for further
emissions from other energy levels. High power is also possible in the MIR with rare-
earth ions. For example, broad band emission around 2 μm with Multi-Watt output
power was achieved in a Tm3+-doped fibre amplifier (TDFA) [77]. A 24 W liquid-
cooled CW 3 μm Er3+-doped ZBLAN fibre laser has also been developed [78]. ASE
waveguide based sources are clearly integratable, the challenge is how to span the wide
range required for sensing of a general nature, though they could be useful for sensing
in specific wavebands, for example for isotopic detection of methane.
Supercontinuum (SC) sources generated from dispersion engineered chalcogenide
fibres or waveguides are able to offer continuous broadband emission in the MIR range
10
with reasonable average power. For instance, broadband supercontinuum spanning from
1.8 μm to >7.5 μm with total power ~20 mW and source brightness > ×100 that of
current synchrotrons was reported from a chalcogenide waveguide excited with ~320 fs
pulses at 4 µm [32]. A supercontinuum with broader band extending from ~2 µm to >10
µm generated using a chalcogenide buried rib waveguide pumped with 330
femtosecond pulses at 4.184 μm was reported by the same group [34]. By launching
ultra-short, high peak power pulses with centre wavelengths of 4.5 μm and 6.3 μm
respectively, into a short piece of chalcogenide glass optical fibre, SC spanning 1.5-11.7
μm and 1.4-13.3 μm were generated, respectively, though at low average powers [30].
Also, MIR supercontinuum generation spanning from ~1.8 to ~10 μm within a dynamic
range of 15 dB has been demonstrated from a 11-cm-long step-index chalcogenide fibre
pumped with ~330 fs pulses at 4.0 μm from an optical parametric amplifier (OPA) [33].
Broadband high spatial coherence sources with wide emission range like this are clearly
highly desirable for molecular sensing, and the potential for integration is clear.
However, to expand the SC emission deep into the MIR, a mode locked MIR pump
source is optimal as shown in the experiments cited above. Further, this needs to
ultimately be integrated to control costs. Fortunately, MIR laser sources are possible
with rare-earth ion doped waveguides in suitable hosts, and so the promise of an
integrated pump source is real. The enormous number of transition lines in different
rare-earth ions cover almost the whole MIR range from 2 to 10 µm, which potentially
offer great support to SC generation as pump sources [35, 39].
(b) Free-space optical communications
Free-Space Optical Communication (FSO) systems are used for high rate
communication between two fixed points over distances up to several tens of
kilometres. In comparison to radio-frequency (RF) counterparts, FSO systems have high
bandwidth, allowing much higher data rates. Besides applications such as, metropolitan
area network (MAN) extension, local area network (LAN)-to-LAN connectivity, high
definition television (TV) and medical image/video transmission, wireless video
surveillance monitoring, FSO systems are considered as an efficient solution for the
“last mile” problem to bridge the gap between the end user and the fibre optic
infrastructure already in place. Currently, FSO link reliability is the biggest challenge
preventing widespread adoption, especially at long ranges due to atmospheric
turbulence-induced fading and sensitivity to weather conditions [79]. The advantages of
MIR range sources for FSO communications arise from the usual λ-4 light scattering
11
dependence (λ is wavelength), making the long wavelength region considerably more
reliable in the dusty and smoky environments, characteristic of many urban areas and
even battlefields [80]. The MIR between 3 to 5 μm is also considered a low loss
atmospheric window and eye safe. Along with the ultra-high bandwidth and highly
secure data transfer for both short and medium range, free-space optical
communications at 3 to 5 μm has been seriously considered as a solution to the “last
mile” communications problem. Thus, besides QCLs, emissions from rare-earth ion
doped materials at 3 to 5 μm wavelengths are in great demand to fulfil this task.
(c) Medical laser surgery
Significant benefits are known to be possible with laser surgery compared with
conventional mechanical cutting methods. For example, laser energy can seal small
blood vessels, so using a laser enables dry field surgery. Also clear field of view, precise
cutting control, and operating in a small area under a microscope could be realised with
laser surgery. Furthermore, the use of femtosecond MIR pulsed sources to ablate tissue
without heating has been shown to result in less tissue damage and better, faster healing
[81-84]. There is also the opportunity to use wavelength specific reactions to target the
specific biological tissue without damaging the surrounding tissues [85].
In medical laser surgery, strong absorption of energy in tissues occurs in both the
ultraviolet and MIR wavelength regions, but because of the potential mutagenic effects
of ultraviolet wavelengths, MIR lasers are preferred for many types of tissue [86].
Surgical lasers have been evaluated by means of achieving controlled removal of tissue
with minimal collateral thermal injury. While water is the dominant chromophore in
tissue, lasers operating at 2.1, 2.94, 6.1, 6.45 and 10.6 μm associated with strong water
absorption are of greatest interest due to their ability to remove tissue effectively with
little concomitant thermal damage. Among these, erbium ion doped lasers at 2.9 µm
(exploiting the water OH-stretch mode) were used leading to a more efficient ablation
of cartilage with significantly less thermal injury [87].
Alternative laser sources in the 6-8 μm range are in need, and pulsed lasers with
proper pulse parameters are also of interest on account of the low thermal damage.
Currently, options for surgical lasers are limited due to their continuous-wave operation
and energy limitations [88-90]. This leaves open an opportunity for rare-earth ion doped
chalcogenide IR laser sources. Low phonon energy and long wavelength transparency
make rare-earth ion doped chalcogenide hosts a promising candidate in MIR laser
12
sources, whilst the large gain bandwidths of rare-earth ion doped chalcogenide host also
bring advantages in realising ultra-short fs pulsed laser sources.
(d) Defence
MIR technologies are used extensively for military purposes, including target
acquisition, surveillance, night vision, homing and tracking. And also some high power
MIR lasers, for example the Mid-Infrared Advanced Chemical Laser (MIRACL), were
developed as directed energy weapons [91, 92]. Compared with lasers in the visible/NIR
wavelength, MIR operation has immediate potential for target illumination and
designation applications, because the visible/NIR wavelength based illuminators and
designators have become less effective due to the availability of night vision sensors
that can be obtained inexpensively and easily by adversaries [93].
Also, MIR technologies will benefit battlefield communication. Currently battlefield
communications are dominated by radio-frequency technologies, but these technologies
are confronted with issues such as simple eavesdropping, jamming and antiradiation
munitions [94]. NIR free-space optical links are relatively immune to these issues but
are degraded by fog and smoke, making it less dependable. Communication in MIR
region is scattered much less by particulates even in fog and smoky conditions, enable
reliable communication in such conditions [95].
In addition, the MIR laser wavelengths are considerably more eye safe than the NIR
laser sources at the same power, because the MIR wavelengths are strongly absorbed by
water and thus cannot reach the retina.
1.3 Rare-earth ion doped chalcogenide glasses
1.3.1 Introduction
Amongst all rare-earth ion elements, erbium ion (Er3+) doped materials have attracted
considerable attention because of their potential applications in optoelectronics, and
their importance in the telecommunication area due to the 1.55 μm emission located at
the low attenuation range of massively deployed silica-based fibre optics [96].
The 1.55 μm emission arising from the 4I13/2→4I15/2 transition in the trivalent state of
erbium ion has been widely used to create optical fibre and waveguide amplifiers in
oxide glass hosts [97-100]. There are some major differences between Er3+ doped
chalcogenide glasses and Er3+ doped oxide glasses, resulting mainly from the low
13
phonon energy of chalcogenide glasses. Compared with the oxides, the low
characteristic phonon energy of the chalcogenide glasses hugely lowers the multi-
phonon processes such as multi-phonon relaxation (MPR) or phonon-assisted parasitic
transitions. This results in a range of processes being present in chalcogenide glasses
that are not available in other hosts, for example, long lived states away from the
traditional 1550 nm erbium (III) transition band allows both extra transitions not seen in
oxides (typically but not exclusively in the MIR) and a range of up conversion, cross-
relaxation, and excited state absorption based processes which have far higher
efficiency than in oxide hosts. Some chalcogenide glasses, such as Ga-containing
chalcogenide or chalcohalide glass [101-104], can also accept high doping levels (1-2
at%) without clustering and this makes it feasible to achieve high gain in a relatively
short device. Additionally, the non-equilibrium growth of thin films via methods such
as co-evaporation or co-sputtering allows the possibility of highly doped films in hosts
where clustering prohibits the formation of useful bulk glasses. Details on options for
film deposition and waveguide fabrication can be found in Section 3.2.3 and 4.4,
respectively. Hence it should, in principle, be possible to use optical gain to compensate
the optical losses in a chalcogenide waveguide as a route to improving the efficiency of
all-optical processing, and to realise MIR waveguide based laser or amplifier devices.
1.3.2 History of rare-earth ion based chalcogenide laser and amplifier devices
There have by now been many reports on the luminescence properties of erbium ion
doped chalcogenide glasses and a smaller amount of work on thin films and
waveguides. At the commencement of this thesis, however, there had been only three
demonstrations of fibre or waveguide amplifiers and a further two bulk glass based
devices using rare-earth ion doped chalcogenide glasses, in spite of the good progress
that has been made on rare-earth ion doped bulk glasses. The first laser action in a rare-
earth ion doped chalcogenide glass was demonstrated by Schweizer et al. in 1996 [105].
In their experiment, CW lasing at 1.08 μm was achieved through a 1.42 mm thick
neodymium-doped gallium lanthanum sulfide (GLS) glass with a concentration of 1.5
mol% (2.6x1020 ions/cm3) Nd2S3 when pumped with a Ti:sapphire laser either at 0.815
or 0.890 µm. Laser action ceased at high pump powers, due to thermal lensing effect. In
the next year, laser action at 1.080 µm was obtained in a 22 mm long GLS glass fibre
under excitation at 0.815 µm [106]. The multimode GLS fibre with a 14 µm core doped
with 0.05 mol% Nd2S3, was fabricated using the rod-in-tube method. Self-pulsing
14
behaviour was noticed in the experiment but the cause for this was unknown. In 2000,
optical amplification at 1.34 µm with a gain coefficient of 0.81 dB/mW was achieved in
a single-mode Pr3+ doped Ga-Na-S (GNS) fibre [107]. The fibre, doped with 750 ppm
Pr3+, had an attenuation loss of 1.2 dB/m at 1.31 µm was fabricated using the extrusion
method. By pumping a 6.1 m long fibre at 1.017 µm, a net gain of 32 dB was achieved
when the pumping power was 90 mW.
In 2002, Mairaj et al. reported lasing at 1075 nm from a neodymium ion doped Ga-
La-S glass channel waveguide written using a focused UV-laser beam (λ=244 nm)
[108]. The maximum measured refractive index change after treatment was Δn~+10-3. A
lasing experiment was performed on a 0.5 mol% Nd2S3 doped 16 mm long waveguide
with propagation loss estimated 0.5 dB/cm. Single mode laser operation at output slope
efficiency of 17% with respect to absorbed power was obtained.
Laser action also has been demonstrated in neodymium ion doped GLS glass
microspheres (1.5 mol% Nd3+ doped) [109]. Chalcogenide spheres of ~100 µm diameter
were fabricated, and the Q factor from this sphere was of the order of 104. Under
excitation at 808 nm, single and multimode laser action was demonstrated at
wavelengths between 1075 and 1086 nm with a pumping threshold of 82 mW. Resonant
peaks shift to longer wavelengths with the increasing of excitation power was observed,
and both thermal and third order nonlinear effects were thought responsible for this
resonant wavelength shift.
Since these demonstrations there has been no gain or lasing demonstrated in
chalcogenide hosts at wavelengths of 1.40 μm or longer. High power MIR amplified
spontaneous emission from rare-earth ion doped chalcogenide fibres was reported by
the Naval Research Laboratory [40, 76, 110]. In their experiment, a 150 μm core, 250
μm cladding Pr3+-doped selenide fibre was pumped using a 1.97 μm diode laser. ASE
emission from 3.2-5.8 μm was achieved. Rise and fall times of the emission from this
fibre source were also studied and it was claimed that the temporal response could be
tailored by selectively modifying the composition of the fibre material.
In addition to rare-earth ions, transition metal (TM2+) ion doped polycrystalline
chalcogenide hosts also offer properties such as ultra-broadband gain, low saturation
intensities and large pump absorption coefficients. This has, until now made them the
gain media of choice for cost effective broadly tunable lasing in the shorter waveband of
the MIR. Amongst the family of TM ion doped chalcogenide glasses, Cr2+:ZnS/Se and
Fe2+:ZnS/Se are the best performers on account of their better thermal properties. With
15
Cr2+ and Fe2+ doped chalcogenide materials, the 1.9-6 μm spectral range is potentially
accessible, though with some limitations/compromises. Lasers with high output powers
(average power of 18 W in gain switched lasing [111] and 30 W in pure CW at 2.4 μm
[112]), large tunable range (over 1800-3100 nm [113]), short-pulse (~40 fs from a
graphene passive mode-locked laser with 250 mW average output power at 2.4 μm
[114]), multi-Joule long-pulse output energy (2.1 J output energy from a low
temperature pulsed Fe2+:ZnSe laser at wavelength of 4.1 μm [115]), and narrow spectral
linewidth (<100 kHz with spectral region from 2.12 to 2.58 μm [116]) were also
realised with TM ion doped chalcogenide glasses[112]. Therefore, it is reasonable to say
that currently this is certainly one of the most effective routes for room temperature
lasing in the 1.9-3 µm and perhaps even ultimately the 1.9–5 µm spectral range though
at minimum thermoelectric cooling is required to access the longer wavelength range.
New schemes for gain element thermal management due to the high potential for
thermal lensing are the key to improve output power in the future and improve the
practical usability of such sources, especially the Fe2+:ZnSe which have to operate at
low temperatures (<-80 °C).
1.3.3 Rare-earth ion doped chalcogenide materials
Amongst chalcogenide hosts, there has been a particular focus has on those containing
Ga, such as Ga-La-S (GLS), Ge-Ga-S and Ge-Ga-Se [46, 117-123], since they have
been shown to accept high Er3+ concentrations without clustering. For example,
Schweizer et al. [117] studied the properties of the Er3+ doped GLS glass system for
MIR applications. Tonchev et al.[118], reported 1550 nm photoluminescence (PL)
decay lifetime around 1.5-2 ms with 975 nm pumping in (As2Se3)1-x(GaSe)x (with Ga
from 0 to 5 at%) bulk glass doped with Er2S3 (1% Er3+). They found that the PL decay
lifetime increased linearly with Ga concentration from about 1.5 ms to 2 ms with the
addition of Ga up to 5%. Allen et al. [119] studied the photoluminescence
characteristics of a series of Er3+ doped chalcogenide glasses and found that all samples
exhibited lifetimes in the 1–4 ms range. The Ga-Ge-As-Se glass had the shortest
lifetimes of 1–1.5 ms for 980 nm pumping, whilst Ga-Ge-Se and Ga-Ge-S samples had
the highest values of 2–4 ms. A strong correlation between the Er3+ ion and Ga
concentrations that affects the properties of Er3+ doped Ge-Ga-Se glasses was found in
[120]. Kasap et al. examined the optical and photoluminescence properties of Er3+
doped Ge-Ga-S glasses of near stoichiometric compositions (Ge28Ga6.2S65.3:Er0.5 at%)
and calculated a lifetime around 2.6 ms using the Judd–Ofelt theory [121]. However,
16
Ga-containing bulk glasses pose significant problems during the thin film deposition
required to make waveguides. In particular, upon melting these glasses tend to phase
separate into a non-volatile Ga-S(Se) phase leading to significant deviations from the
starting composition in the film, typically the film being gallium poor. As a result, to the
best of our knowledge there have been no reports of stable, high quality thin films or
waveguides (loss of less than 1 dB/cm) made from Ga-containing chalcogenide glasses.
Chalcohalide glasses are also reported with high rare-earth solubility and promising PL
emissions in MIR [104, 124, 125], however, due to the lack of experience in
chalcohalide film deposition and waveguide fabrication in our group, no research on
chalcohalide was conducted in this thesis.
There have, however, been many demonstrations of high quality chalcogenide thin
films and waveguides in other chalcogenide hosts such as As2S3, Ge-Sb-S and
Ge11.5As24Se64.5 (at%) [4, 11, 13, 14]. For example, loss values as low as 0.05 dB/cm at
1550 nm have been measured for un-doped 4 μm wide As2S3 rib waveguides etched into
2.5 μm thick films, whilst losses for highly nonlinear dispersion-engineered As2S3
waveguides with a nonlinearity, γ≈10 W−1m−1 are now as low as 0.3 dB/cm [13].
Recently, Ge11.5As24Se64.5 (at%) nanowires (585×575 nm) with extreme nonlinearity
coefficient of γ=130 W−1m−1 and moderate losses of ≈1.65 dB/cm were also reported
[65]. Propagation loss as low as 0.42 dB/cm was demonstrated in submicron Ge23Sb7S70
(at%) waveguides (700×600 nm ribs with 290 nm etch depth) using chlorine plasma
etching [14]. These materials, have to date, proven to be the most suitable for planar
waveguides [4, 11]. However, doping of rare-earth ions into bulk glasses with these
compositions is difficult because the absence of Ga means the solubility of rare-earth
ions in such glasses is low. As2S3 bulk glass doped with Er2S3 to give 0.1 at% Er
concentration produced complex narrow line structures in the PL spectrum that are
similar to those found in crystalline material such as Er3+:YAG [126], implying low
erbium ion solubility in As2S3 bulk glass; while for the ternary Ge33As12Se55 (at%)
glass, erbium ion clustering was observed at concentrations above 0.2 wt% [127].
Unfortunately, useful waveguide based devices require ~1-5 at% doping levels in order
to obtain the ~1dB/cm gain required for practical devices.
Interestingly, there is strong evidence that some films can incorporate larger
amounts of rare-earth ions compared with bulk samples when prepared by physical
vapour deposition methods. This is most probably due to the fact that films are created
in strongly non-equilibrium conditions by condensing a vapour onto a cold substrate.
17
This means that single isolated rare-earth atoms (or ions) are immediately incorporated
into the film thereby inhibiting clustering. This contrasts sharply with the situation used
to create a bulk glass where the dopant has to be soluble in the molten host. Lyubin et
al. [128] reported that co-thermal evaporation of erbium with As2S3 produced films with
Er3+ concentration as high as 4 at% without any signs of clustering and led to strong PL
emission under Ar+ laser excitation at 514 nm. Vigreux-Bercovici et al. [129] also
reported sputtering of a 3% Er3+:As2S3 composite target to produce a thin film that had
1.5 μm transition lifetime of 4 ms. Whilst this is encouraging, the problem of etching
rare-earth ions doped films remains and hence it is sometimes attractive to introduce the
dopant after waveguide fabrication using ion implantation. In this approach a
chalcogenide host that makes the most stable low loss waveguides can be employed.
Fick et al. observed a strong Er3+ emission at 1.54 μm from erbium ion implanted As2S3
and As24S38Se38 (at%) films with a lifetime of 2.3 ms [130]. In Ivanova’s work [131],
the PL properties of Ge-S-Ga films ion implanted with relatively low energy (320 keV)
ions at different fluences was investigated. They reported that the PL efficiency reduced
with increasing Er3+ concentration and that thermal annealing at 230 °C approximately
doubled the PL efficiency at all doses.
Whilst rare-earth ion doped chalcogenide bulk glasses have been intensively
investigated, e.g. [119, 120, 122, 127, 132], as noted above, at the outset of this work
there have been only five demonstrations of gain or lasing in chalcogenide glasses.
Despite the many rare-earth elements added in bulk glass and the rich vein of possible
transitions at longer wavelengths, there have since been no further demonstrations of
amplifiers or lasers. A brief collation of representative measured NIR and MIR
emissions of various rare-earth ion doped chalcogenide glasses and fibres are shown in
Table 1.3.
18
Table 1.3 Collation of the MIR emission of rare-earth ion doped chalcogenide glasses and
fibres.
Dopant(s) Host glass Excitation
/ µm
Emission / µm Transition Reference
Er3+ As2S3 1.48 1.54 4I13/2→4I15/2 [133]
GLS 0.66/0.81 2.0/2.75/3.6/4.5 [117]
GeGaSbS 0.804 4.3-4.8 4I9/2→4I11/2 [134]
Ho3+ GeGaAsS/Se 0.9 1.6 5I5→5I7 [135]
GeAsS 0.905 1.2/2.0/2.9 5I6→5I8 /5I7→5I8
/5I6→5I7
[136]
GLS 0.76/0.9/2 1.2/1.25/1.67/2/2.2/2.
9/3.9/4.9
[137]
Tm3+ GeAsS 0.698/0.8 1.2/1.4/1.8 1H5→3H6
/3H4→3F4
/3F4→3H6
[138]
GeGaS-CsI 0.8 1.48/1.8/2.3/2.8 [139]
GLS 0.7/2 3.8/5.38 3H5→3F4 /
3F2,3→3H4
[48]
Dy3+ GeAsS
GeGaS
0.808 1.33/1.75/2.9/4.38/5.
27 (No spectrum)
[140]
GeGaS 0.798 2.9 (Tm3+/Dy3+ co-
doped)
6H13/2→6H15/2 [141]
GaSbS 1.32 2.95/3.59/4.17/4.4 [142]
‘selenide’ ~4.5 6H11/2→6H13/2 [103]
GeAsGaSe 1.3 3.0/3.2/4.5/5.5/7.6 [54]
Pr3+ ‘selenide’ 3.5-5.5 [41]
GeAsGaSe 2 3.4/4.0/4.8/5.2/7 [54]
Tb3+ GeAsGaSe 1.97 3.1/4.7/4.8/7.5/10.5 [54]
GLS 2 4.8/8.1(No spectrum) 7F5→7F6 /7F4→3F5 [48]
19
1.3.4 Potential issues for rare-earth ion doped chalcogenide glasses
Whilst chalcogenide glasses have much promise as outlined above, there are also
potential issues that need to be outlined and considered in the forthcoming work. For
example, it is well known that in erbium ion doped silicon, photoluminescence can be
excited and observed via the energy transfer from an electron-hole pair trapped at an
erbium-related defect state in the silicon band gap. In detailed studies, the reverse
process was also observed. In the erbium ion doped silicon system, energy transfer from
excited erbium ion back to the electronic system of the silicon host was observed and
named energy back-transfer. Obviously, this process is detrimental for devices such as
amplifiers or lasers, for it depopulates the excited ions in a non-radiative way and
causes thermal degradation.
A similar phenomenon has been noticed in rare-earth ion doped chalcogenide
glasses. In Gu et al.’s study [143], careful investigation via PL and photoluminescence
excitation spectroscopy (PLE) in Er3+ doped As2S3 and Ge33As12Se55 (at%) glasses,
remarkably broad PLE spectra for the 1550 nm emission extending from the Urbach
absorption edge to beyond 1000 nm were observed [143, 144]. Similar features were
also observed in the PLE spectrum of a 1-μm-thick Er3+-doped As40Ge10Se25S25 (at%)
film deposited on silicon substrates using radio frequency (RF) sputtering [145]. This
feature indicated that, like what happened in rare-earth ion doped silicon, there was
another energy transfer mechanism by which energy can be directly transferred in either
direction between the chalcogenide host and rare-earth ions. Whilst this energy transfer
mechanism potentially enables flexible pumping wavelengths and more distantly the
possibility of direct electrical pumping, the back-transfer process could also prevent
rare-earth ion doped devices from attaining gain.
Alternately, Ivanova et al. examined the conductivity properties of erbium ion doped
Ga-Ge-S-Se glasses, and also estimated the efficiency of the energy back transfer
process to the chalcogenide matrix by careful lifetime measurement in erbium ion doped
Ga-Ge-S glasses [146-148]. Samples were exposed to sub-bandgap light, the un-doped
sample showing no change in conductivity; while the sample doped with erbium had a
significant increase in conductivity. It was believed that the erbium was excited to an
energy state above the band gap by a multi-photon mechanism and then interacted with
electrons in the glass matrix to generate additional free electrons thereby increasing the
conductivity. In conclusion, it was pointed out: energy transfers from an excited erbium
20
ion to the host (energy back-transfer) may also be possible. This obviously brings an
obstacle to realise an amplifier or laser based on rare-earth ion doped chalcogenide
glasses.
Ion-to-ion interaction is another issue to be confronted in rare-earth ion doped
materials. With high rare-earth dopant concentrations, the distance between rare-earth
ions becomes small enough that electric dipole-dipole interactions between different
rare-earth ions takes place. The strength of this interaction is strongly related to the
distance between the related ions with a 1/r6 relationship, where r is the inter-ion
distance [149]. The occurrence of this phenomenon can drop the fraction of excited
rare-earth ions significantly at a given pump power, which will in turn degrade the
performance of the device. Publications concerning these phenomena have been
reported in variety of host materials, while one representative of chalcogenide glasses,
sulfide glasses, was claimed to suffer more severe ion-to-ion interaction [150]. In that
work, both Er3+/Ce3+ co-doped tellurite and sulfide glasses were prepared. Up-
conversion emission at 816 nm increased and the lifetime of the 4I13/2 state decreased
with increasing cerium concentration, implying the occurrence of the
4I13/2:4I13/2→
4I15/2:4I9/2 process. While similar phenomena were observed in the tellurite
glasses, they were much weaker, indicating that sulfide glasses suffered more serious
ion-to-ion interactions, which even quenched the 4I13/2 state population. Thus, ion-to-ion
interaction was thought to be another issue preventing researchers from achieving gain
in Er and potentially other doped chalcogenide devices.
Impurity multiphonon relaxation is suggested as an additional method of non-
radiative depopulation of the excited state in the MIR, besides lattice multiphonon
relaxation [39]. In gallium-lanthanum-sulfide (GLS) glass, large oxide additions
(La2O3) increase glass stability to avoid crystallization during fibre drawing. With the
appearance of La2O3, a mixed sulfur-oxygen species is formed leading to a vibrational
absorption at 8.6 µm (1163 cm-1), and this value is thought a more realistic phonon
energy in this oxysulfide glass instead of the value of ~425 cm-1. With this number, only
4 phonons are required to prevent the ~ 2 µm emission, thus it is easy to explain that
GLS glasses have failed to lase in MIR. In more covalent chalcogenide glasses such as
Ge-As-Se, Ge-Sb-Se, oxide exists as an impurity manifesting extrinsic absorption bands
in the infrared spectra of glasses. For example, As-O and Ge-O impurity have
vibrational absorption band centred at ~7.9 µm (~1266 cm-1). These extrinsic impurities
close to a rare-earth ion in a glass host lattice could cause impurity multiphonon
21
relaxation of the rare-earth ion excited states in an analogous way to lattice multiphonon
relaxation by means of local independent oscillators. In this scenario, the probability of
non-radiative emission by impurity multiphonon decay matters. Oxide impurities at low
level of 200 ppmw had great impact on the photoluminescence spectrum in Pr3+-doped
GeS2-based glasses was also reported in [151]. Of course the magnitude of this effect is
directly proportional to the impurity concentration and it does not affect all the rare-
earth ions in the material, but it can have sufficient impact to prevent gain being
demonstrated. In Section 5.2 of this thesis, the PL of RF sputtered Er3+:As2S3 films
drops almost down to zero in some conditions, which is believed due to the existence of
O-H impurities which are resonant in energy with the desired transition. This has also
been documented in tellurite glass planar waveguide amplifiers as a mechanism
preventing gain [152]. Thus in any investigation into rare-earth ion doped gain, material
optimisation is necessary to rule out contaminant species as significant contributors to
non-radiative decay.
In the decades following the initial demonstration of the chalcogenide laser, and
despite considerable efforts, no successful result beyond 1300 nm has been reported.
With the huge potential applications in MIR sensing, and all optical signal processing
and telecommunications, erbium ion doped chalcogenide amplifiers and lasers were
intensively studied, but no significant results have been achieved. In addition to this, to
date only one explanation has been proffered as to why rare-earth ion doped
chalcogenide glasses do not perform as expected. This is due to absorption and non-
radiative decay of excited rare-earth ion by extrinsic phonons due to impurities in the
chalcogenide glass host as mentioned above [39].
1.4 Problem under study
As has been laid out above, since the first demonstration of a laser using Nd3+ at 1080
nm and Pr3+ at 1300 nm in both bulk glass and fibres, no real progress has been made on
other ions or longer wavelengths in chalcogenide hosts. Given the clear MIR
capabilities of chalcogenide glasses and their potential as all optical processing hosts at
1550 nm, this clearly needs to be investigated. Likewise, planar waveguide based lasing
devices have not prospered despite some promising looking initial results in thin films.
In neither case has much been said in the literature about the reasons why progress has
not been achieved, apart from the reason of extrinsic non-radiative decay [39].
22
Therefore, the plan for this thesis was to investigate these two issues. The ion of
choice for the studies was erbium, as it provides gain at 1.5 µm for optical signal
processing applications, as well as MIR transitions at 2.9, 3.4 µm and 4.5 µm. In
addition, it has potentially efficient multiple pump photon upconversion to the visible
where the interactions with the band gap may be studied. The objective of this study
was to make working rare-earth ion doped chalcogenide waveguide amplifiers and
lasers at 1.5 µm, and potentially demonstrate MIR amplification in a planar format.
1.5 Outline of this work
This thesis studies both the materials and waveguide structures for rare-earth ion doped
chalcogenide glasses. Several deposition modalities are covered, comparisons made to
bulk glasses, and waveguide devices fabricated and tested. The remainder of this thesis
is organized as follows:
Chapter 2 introduces comprehensively the important properties of rare-earth ions,
focussing especially on erbium ion and characterisation methods of rare-earth ions in
hosts. Factors related with lifetime and emission properties are discussed. Important
processes in rare-earth ion doped glasses, such as energy migration processes,
concentration quenching and solubility of rare-earth ions in chalcogenide hosts are also
discussed.
Chapter 3 focuses on As2S3, the main workhorse of this thesis. Optical and
structure properties such as transmission, absorption, dispersion and Raman spectrum of
As2S3 are introduced at the first part. Bulk As2S3 was fabricated by melt-quenching
method, and then erbium ion doped As2S3 films were deposited using co-thermal
evaporation. The physical and optical properties of the obtained films are detailed. The
influence of thermal and light post-treatment and even higher temperature (above the
transition temperature of As2S3) post-treatment on the emission properties of erbium ion
are studied. Cooperate up-conversion phenomena in erbium ion doped As2S3 film were
observed and investigated as well. With all this information, an optimised erbium ion
concentration was arrived at, and this concentration was applied in the following
chapter.
Chapter 4 covers the basic theory of light propagation in waveguide, waveguide
design and simulations. Performance of erbium ion doped As2S3 waveguide amplifier
was also simulated using the OptiSystem commercial software, and promising results
were achieved. Several methods of waveguides fabrication are detailed. Photo-
23
lithography combined with inductively coupled plasma reactive ion etching was applied
in this work to pattern erbium ion doped As2S3 waveguides. Erbium ion doped As2S3
waveguides with a low loss of 0.35 dB/cm were achieved, and for the first time internal
gain from 1570 to 1630 nm was observed in an erbium ion doped chalcogenide
waveguide amplifier. Neodymium ion doped As2S3 waveguides were also fabricated
and tested, but no gain/lasing was observed. Given the test results and information from
Raman spectroscopy, neodymium ion clustering in the As2S3 host was thought the main
reason for failure.
Chapter 5 introduces different approaches to realise better population inversion in
erbium ion doped As2S3 waveguides. ErCl3 and Er2S3 were used in film deposition
instead of erbium metal, however in both cases film quality became an issue that
prevented further progress. RF sputtering of As2S3 together with a piece of erbium foil
was also performed. The PL intensity from obtained sputtered films (after thermal post-
treatment) was even stronger than the PL intensity from the film on which internal gain
was achieved. However, due to the columnar growth habit RF sputtering has, the O-H
contained in the column-structure quenched PL rapidly, so no amplification result was
achieved from this waveguide.
Chapter 6 reports the research results on erbium ion doped chalcogenide glasses
and waveguide amplifiers based on Ge-Ga-Se hosts, which are known to have better
rare-earth ion solubility than As2S3. A series of erbium ion doped Ge-Ga-Se glasses was
synthesised, and their optical properties are measured and analysed. Waveguide
structures were patterned on erbium ion doped Ge-Ga-Se films deposited using co-
thermal evaporation, and the optical properties were characterized. The film deposition
equipment was modified to overcome the particulates on deposited film, after that high
quality films were finally deposited. With these high quality films, erbium ion
population inversion up to 50% was achieved, but appearance of photo-darkening
phenomenon was observed in amplification experiments which presented a hurdle to
realising a working rare-earth ion doped chalcogenide waveguide amplifier in practice.
Chapter 7 summarises the main challenges, achievements of the work then suggests
future work to further advance the field of rare-earth ion doped chalcogenide waveguide
devices.
24
25
Chapter 2
Rare-earth ions in chalcogenide glasses
2.1 Introduction to rare-earth ions
The rare-earth elements (also referred to as the lanthanides) are a set of seventeen
elements in the Periodic Table, specifically the fifteen lanthanides, as well as scandium
and yttrium. In the lanthanide series, the 4f-shell is consecutively filled from cerium
(4f15d16s2) to lutetium (4f145d16s2). Compared with the 4f-electons, the more distant
5d1- and 6s2-electrons with lower ionization energies can be removed much more easily
and this explains the similar reactivates that the lanthanides share. After removing the
5d1- and 6s2-electrons, the 4f-electrons are still shielded by 5s- and 5p-electrons,
therefore, they are relatively insensitive to the host environment. The 4f electron shell
has a series of energy levels which are annotated by the orbital angular momentum
number L, the overall spin S, and the total angular moment J. These levels are known as
the LSJ levels, and written as 2s+1Lj. In a neutral environment, these multiple energy
levels are degenerate and electric dipole transitions between different levels are
forbidden. But if rare-earth ions are subjected to the crystal field of a host matrix, then
despite the 5s- and 5p-electron shielding, there is still sufficient interaction between the
4f-electrons and the crystal field to affect the energy levels [153]. As a result, a series of
discrete energy levels are created, and a radiative electric dipole transition between
different multiplets is then allowed to occur. Because of the weak oscillator strength of
the electric dipole transition under these circumstances, narrow emission lines are
observed in crystalline hosts for all the rare-earth ions, and the emission wavelengths
are only slightly influenced by individual ligand systems. One crucial advantage
supporting rare-earth ion luminescence is the long lifetimes of their excited states which
are directly correlated to the low probabilities of forbidden transitions. As an example of
interest in this research, Figure 2.1 shows the effect of spin-orbit and crystal field
interactions on the energy levels of erbium ions in a solid host [154]. All the possible
inter-level emission wavelengths are calculated based on the data from single crystal
erbium oxide [155], and the detailed results are shown in Table 2.1.
26
Figure 2.1 Energy levels of Er3+ quantum energy with effect of spin-orbit and crystal
splitting and excitation wavelengths from 4I15/2, 4I13/2 and 4I11/2 to various levels [154], see
Table 2.1
Table 2.1 Calculated possible inter-level emission wavelengths in µm with data of Er2O3
crystal [155].
4I15/2 4I13/2 4I11/2 4I9/2 4F9/2 4S3/2 2H11/2
4I13/2 1.548
4I11/2 0.992 2.759
4I9/2 0.815 1.719 4.562
4F9/2 0.664 1.162 2.008 3.587
4S3/2 0.554 0.862 1.253 1.728 3.333
2H11/2 0.528 0.802 1.131 1.504 2.589 11.59
4F7/2 0.494 0.725 0.983 1.253 1.926 4.56 7.52
4F5/2 0.457 0.648 0.847 1.04 1.465 2.61 3.37
4F3/2 0.499 0.632 0.82 0.999 1.385 2.37 2.98
2H9/2 0.412 0.561 0.704 0.832 1.083 1.6 1.86
27
For Er3+, the 4I13/2 energy level is treated as the metastable excited state as it has
essentially the longest excited state lifetime of the accessible levels in most hosts. The
radiative lifetime decay from this level to the ground state 4I15/2 is relatively long, up to
tens of milliseconds in some host materials, though in low phonon energy materials
such as tellurite and chalcogenide glasses, the 4I11/2 and 4I9/2 state lifetimes become
comparable to the lifetime of the 4I13/2 state [46]. Therefore, several different pumping
schemes were proposed. Typically, in higher phonon energy hosts such as silica, an
excitation laser at 980 nm is used in a three energy level system. Ions are excited from
the ground state 4I15/2 to the excited state 4I11/2, and then decay to a metastable state 4I13/2
quickly due to the interaction between the excited erbium ions and the thermal phonons
in the host material. As erbium ions accumulate in the long lived 4I13/2 energy level,
population inversion and therefore optical gain is achievable for this scheme. Pumping
with 1490 nm is another option. In this scheme, erbium ions are excited into the top of
the 4I13/2 manifold directly by absorbing pump photons. These excited ions then de-
excite down through the sublevels which are broadened to the point of strong overlap by
thermal phonons to form a population inversion between the lower energy part of the
4I13/2 manifold and the ground state 4I15/2. In this case, the energy level 4I13/2 is
considered to be effectively split into two parts that play different roles in the process.
Therefore, in this pumping scheme, the erbium ions are treated as a quasi-three energy
level system. However, as the pump photon energies lie inside the excited state
manifold, there is some pump stimulated emission and this places limits on the
maximum inversion achievable. It is also possible to pump Er3+ via the even higher
energy levels at 800 nm and 670 nm.
2.2 Properties and characterization of rare-earth ions in
glasses
2.2.1 Lifetime of rare-earth ions in an excited energy level
The lifetime, τ, of an excited energy level is the inverse of the probability per unit time
of the exit of an ion from that excited level. In the ideal case, the decay of the
population in a given level drops exponentially with a time constant equal to the
lifetime. There are a number of pathways for the excited state decays. The total
probability is equal to the sum of the individual pathway probabilities. Each pathway
can, therefore, be considered to have a separate lifetime. Usually they are classified as a
28
radiative lifetime τr or non-radiative lifetime τnr. The total lifetime can then be written
as:
1 𝜏 = 𝑊𝑡𝑜𝑡 = 𝑊𝑟 +𝑊𝑛𝑟⁄ (2.1)
The overall emission probability for an excited energy level depends upon the
relative sizes of Wr (radiative emission) and Wnr (the probability of non-radiative
emission) per unit time, which will be elaborated in the following parts, (a) and (b),
respectively.
(a) Radiative Relaxation
Radiative relaxation refers to the situation that excited ions decay to a lower state by
emitting a photon. When two energy levels, i and j, are sufficiently close they become
thermally linked and their relative populations are expressed by Boltzmann's
distribution law [156, 157]:
𝑁𝑖
𝑁𝑗= 𝑒𝑥𝑝 (−
∆𝐸𝑖𝑗
𝑘𝐵𝑇) (2.2)
where Ni,j are the population of upper and lower states, ΔEij is the energy difference
between state Ni and Nj, kB is the Boltzmann constant and T is the absolute temperature.
Obviously, if ΔEij is big enough, the ratio of the population between the upper and lower
states becomes essentially zero. This means the two energy levels are not thermally
coupled and the relaxation between these two levels can occur with relatively high
probability via photon emission or other non-radiative routes. There are two types of
mechanisms for photon emission, depending on whether population inversion between
these two levels has been formed, they are known as spontaneous or stimulated
emission.
(b) Phonon Relaxation
In reality, non-radiative emission has many contributing components, such as, lattice
multiphonon decay and impurity/defect multiphonon decay processes, expressed
mathematically as:
𝑊𝑡𝑜𝑡 = 𝑊𝑟 +𝑊𝑚𝑝 + ∑𝑊𝑖𝑚𝑝 (2.3)
where Wmp is lattice multiphonon decay rate and ∑𝑊𝑖𝑚𝑝 is the sum of all impurity
multiphonon decay rate [39]. Whilst the exact mechanisms underlying some of these
decay processes can be complex, the model described by Layne [158] has usually been
found to be an acceptable approximation for Wmp. Under this description, if the energy
29
gap to the nearest lower Stark manifold of the rare-earth ions is ΔE, then normally p
phonons from the lattice are needed to bridge this gap, where 𝑝 = ∆𝐸 ℏ𝜔𝑙𝑎𝑡𝑡𝑖𝑐𝑒⁄ . For
energy gaps that can be bridged by a few phonons, the multiphonon rate can be
expressed empirically as follows [39, 159]:
𝑊𝑚𝑝(𝑇) = 𝛽[1 + 𝑛(𝑇)]𝑝𝑒−𝛼∆𝐸 (2.4)
where n(T) is the number of thermally generated phonons per lattice mode at absolute
temperature T, which can be defined through the Bose-Einstein equation as:
𝑛(𝑇) = [𝑒𝑥𝑝(ℏ𝜔 𝑘𝑇⁄ ) − 1]−1 (2.5)
β is a material constant, ΔE is the energy gap between two successive levels, and α is
expressed by
𝛼 = −(ℏ𝜔)−1𝑙𝑛(𝜖) (2.6)
where, ϵ is the coupling constant.
α, β and ϵ are dependent on the host but independent of the specific electronic level
of the rare-earth ion from which the decay occurs. Typical non-radiative parameters for
different glass hosts are shown in Table 2.2.
It is clear for a given glass host, the probability of non-radiative relaxation between
two energy levels depends on the energy gap ΔE and how many phonons are required to
bridge this gap. Amongst all the host materials listed in Table 2.2, chalcogenide glasses
have the lowest minimum phonon energy, and the five to seven orders of magnitude
lower multiplicative β factor. This determines that chalcogenide glasses have the lowset
non-radiative relaxation probability, therefore longer wavelength emissions that could
not be realised in other hosts due to high phonon energy could be expected in
chalcogenide hosts.
30
Table 2.2 Typical non-radiative parameters for different glass hosts (a from [67]; b from
[160]).
Glass β (s-1) α (cm) ħω (cm-1) ϵ
Phosphatea 5.4x1012 4.7x10-3 1200 0.0036
Silicatea 1.4x1012 4.7x10-3 1100 0.0057
Germanatea 3.4x1010 4.9x10-3 900 0.012
Telluritea 6.3x1010 4.7x10-3 700 0.037
Fluoridea 1.88x1010 5.77x10-3 500 0.056
Borateb 2.9x1012 3.80 x10-3 1400 0.0048
Ge-Ga-Sb 8.13x105 2.83 x10-3 350 0.371
Ge-As-Sb 2.56x106 2.95 x10-3 350 0.356
La-Ga(Al)-Sb 1x106 2.9 x10-3 350 0.362
Chalcogenidea 1x106 2.9x10-3 300 0.419
2.2.2 Absorption and emission cross-section
(a) Absorption cross-section
Absorption cross-section can be established by direct light transmission measurements.
To obtain an absorption spectrum, a broadband light source is usually used in
conjunction with a monochromator to provide wavelength discrimination. When a
sample is illuminated with beam of intensity I0(λ), the output power intensity I(L,λ) can
be written according to the Beer-Lambert law:
𝐼(𝐿, 𝜆) = 𝐼0(𝜆)𝑒𝑥𝑝(−𝛼(𝜆)𝐿) (2.7)
where α(λ) is the absorption coefficient and L is the optical path length of the sample
along the propagation direction. The absorption cross-section σa(λ) is then defined as the
absorption coefficient, α(λ), normalised by the ion concentration, N:
𝜎𝑎(𝜆) =𝛼(𝜆)
𝑁=
1
𝑁𝐿𝑙𝑛
𝐼0(𝜆)
𝐼(𝐿,𝜆) (2.8)
31
(b) Emission cross-section
The estimation of the emission cross-section is more complicated than the absorption
cross-section. When it is not straightforward to measure both the absorption and
stimulated emission cross-sections for a rare-earth ion doped material, the Einstein A
and B coefficients for a two-level system are often employed to calculate one from the
other. In a two level system, both lower state (level 1) and upper state (level 2) are split
into multiple components, and the following relationship holds [161]:
𝑔1 ∫𝑣2𝜎𝑎(𝑣)𝑑𝑣 = 𝑔2 ∫𝑣
2𝜎𝑒(𝑣)𝑑𝑣 (2.9)
where 𝑔𝑖 is the degeneracy of level i, v is the photon frequency, σa and σe are absorption
and emission cross-sections, respectively. This equation is a more general form of the
Ladenburg-Fuchtbauer relationship [162], and the validity condition in a rare-earth ion
system is that either the two levels must be equally populated, or all the transitions must
have the same rate. Unfortunately, erbium ion doped glasses cannot meet either of these
conditions. The manifold width of the 4I15/2 and the 4I13/2 states of erbium ion doped
glass are typically a few meV (300-400 cm-1), which is larger than kT=200 cm-1 at room
temperature. In addition, low temperature absorption and emission measurements
indicate the transition strength is sensitive to the Stark levels involved, making it
impossible for all transitions to have the same rate. If an emission spectrum is captured
and the metastable state radiative lifetime is accurately measured, then a formula
derived from the Ladenburg-Fuchtbauer relationship can be applied to calculate the
emission cross-section [97, 161]:
1
𝜏𝑒=
8𝜋𝑛2
𝑐2∫𝑣2𝜎𝑒(𝑣) 𝑑𝑣 = 8𝜋𝑛2𝑐 ∫
𝜎𝑒(𝜆)
𝜆4𝑑𝜆 (2.10)
where τe is the radiative lifetime of metastable state, n is the refractive index of the host
material, c is the light velocity in vacuum, and σe is emission cross-section. This method
is known to work well for erbium ion doped fluorophosphate glass [161].
On the other hand, a relationship between absorption and emission cross-sections
was proposed by McCumber in 1964 [156], as follows:
𝜎𝑒(𝑣) = 𝜎𝑎(𝑣)𝑒𝑥𝑝 (𝜀−ℎ𝑣
𝑘𝑇) (2.11)
where σe is the stimulated emission cross-section, v is the photon frequency, h is Plank’s
constant, k is the Boltzmann constant, and ε is the effective energy required to excite
one Er3+ from the lower state of interest to the upper state of interest at temperature T.
The only assumption required by this theory is the time needed to establish a thermal
32
distribution within each manifold must be shorter than the lifetime of the manifold.
Thus, an accurate value of ε is essential to calculate emission cross-sections.
Miniscalco and Quimby [161] suggested a simple method for approximation of ε by
using the room temperature absorption and emission spectrum half-width. The ion
distribution between the ground level and excited level can be described as follows
[161]:
𝑁1
𝑁2= 𝑒𝑥𝑝(
𝜀
𝑘𝑇) (2.12)
where N1 and N2 are the equilibrium population of ground level and excited level at
temperature T. For Er3+, if all eight Stark components of the ground state 4I15/2 and
seven Stark components of the excited state 4I13/2 are known, the ion distribution can
also be expressed as [161]:
𝑁1
𝑁2=
1
𝑒𝑥𝑝(−𝐸0 𝑘𝑇⁄ )×
1+∑ 𝑒𝑥𝑝(−𝐸1𝑗 𝑘𝑇⁄ )8𝑗=2
1+∑ 𝑒𝑥𝑝(−𝐸2𝑗 𝑘𝑇⁄ )7𝑗=2
(2.13)
where E0 is the separation between the lowest components of each manifold, Eij is the
energy difference between the jth and the lowest component of level i. The electronic
structure could be simplified by assuming the Stark levels for a given manifold are
equally spaced, which means Eij=(j-1)Ei. This reduces the unknown parameters from 14
to 3: E0 and the manifold spacing ΔE1 and ΔE2. From low temperature measurement
results, the absorption and emission peaks corresponding to the transition between the
lowest components of each manifold, therefore, the average value of absorption and
emission peaks are taken as the value of E0. Also the seven ΔE1 spaced levels are chosen
based on the low-energy half-width of the room temperature emission spectrum, while
the six ΔE2 spaced levels are chosen from the high-energy half-width of the room
temperature absorption spectrum accordingly. With these three parameters (E0, ΔE1 and
ΔE2), the N1/N2 value is calculated, then according to equation 2.12, the parameter ε,
which is essential to emission cross-section calculation, could be extracted. The
accuracy of the calculated emission cross-section using this method was verified by
Miniscalco and Quimby in Al/P-silica fibre and fluorophosphate glasses [161]. This
method is also widely applied on rare-earth ion doped chalcogenide glass hosts with
reasonable accuracy [48, 120, 163, 164].
33
(c) Maximum pump efficiency
Knowing the absorption and emission cross-sections, the fraction of ions in the 4I13/2
metastable state can be estimated as a function of excitation wavelength. Under high-
power excitation, the expression for the maximum achievable excited state fraction is
[97]:
𝜂𝑚𝑎𝑥 =𝑛2
𝑛=
𝜎𝑎(𝜆𝑝)
𝜎𝑎(𝜆𝑝)+𝜎𝑒(𝜆𝑝) (2.14)
where n is the total Er3+ concentration, n2 is the population of the 4I13/2 state, σa is the
absorption cross-section, σe is the emission cross-section and λp is the excitation
wavelength.
2.2.3 Judd-Ofelt theory
In order to determine the transition probabilities or oscillator strength of any particular
transitions between energy levels in rare-earth ions, a theory was proposed
independently by Judd and Ofelt in the 1960s [165, 166], since termed the Judd-Ofelt
theory. Since its proposal, it has been widely applied to calculate 4f transition intensities
of rare-earth ions doped in many hosts. This method is based on three Judd-Ofelt
parameters Ω2,4,6 calculated from experimental data normally obtained from ground
state absorption spectra. With these three parameters, important spectroscopic
characteristics like the spontaneous emission probabilities, the radiative lifetimes, the
oscillator strengths, the branching ratios and the quantum efficiency of the levels can be
evaluated.
The relationship between transition oscillator strength and integrated absorption for
a particular transition is given by:
𝑓 =𝑚𝑐
𝜋𝑒2𝑁
9𝑛
(𝑛2+2)2∫𝑘(𝜆)𝑑𝜆 (2.15)
where N is the number density of rare-earth ions in the glass, m and e are the mass and
charge of electron, c is the vacuum velocity of light, n is the refractive index, and k(λ) is
absorption coefficient at wavelength λ, which can be measured from experiment [146].
On the other hand, the oscillator strength of a transition from a level J to a level Jʹ
can be expressed theoretically as follows:
𝑓(𝐽: 𝐽′) =8𝜋2𝑚𝑐
3ℎ(2𝐽+1)𝜆𝑛2× [𝜒𝑒𝑑𝑆𝑒𝑑(𝐽: 𝐽′) + 𝜒𝑚𝑑𝑆𝑚𝑑(𝐽: 𝐽′)] (2.16)
34
where h is the Planck’s constant, 𝑆𝑒𝑑(𝐽: 𝐽′) and 𝑆𝑚𝑑(𝐽: 𝐽′) are the electric and magnetic
dipole transition line strengths, respectively; 𝜒𝑒𝑑 = 𝑛 (𝑛2 + 2)2 9⁄ and 𝜒𝑚𝑑 = 𝑛3 are
the local field correction factors for electric dipole moment transition and for magnetic
dipole moment transition, respectively [167].
The contribution given by electric dipole to the line strength can be expressed as
follows, based on Judd-Ofelt theory:
𝑆𝑒𝑑(𝐽: 𝐽′) = ∑ Ω𝑡|⟨4𝑓𝑁(𝑆𝐿)𝐽||𝑈(𝑡)||4𝑓𝑁(𝑆′𝐿′)𝐽′⟩|
2𝑡=2,4,6 (2.17)
While the line strength resulting from magnetic dipole transitions is expressed by:
𝑆𝑚𝑑(𝐽: 𝐽′) =ℎ2
16𝜋2𝑚2𝑐2|⟨4𝑓𝑁(𝑆𝐿)𝐽| |
𝐿→+
2𝑆→| |4𝑓𝑁(𝑆′𝐿′)𝐽′⟩|
2
(2.18)
where the three terms ⟨‖𝑈(𝑡)‖⟩ are the reduced matrix elements of the unit tensor
operator whose values depend on the rare-earth ion type, and details are given in ref
[168]. The coefficients Ω2,4,6 are intensity parameters which depend on the glass matrix.
By fitting experimentally obtained oscillator strengths with this theoretical expression
via least squares fitting, the Judd-Ofelt parameters Ω2,4,6 can be extracted. These
parameters then allow one to characterise all aspects of emission from the ions in that
host.
As the Ω2,4,6 parameters do not depend on the rare-earth ion transition levels, the
spontaneous emission probabilities between J and Jʹ can be calculated from the
following equation:
𝐴(𝐽′: 𝐽) =64𝜋4𝑒2
3ℎ(2𝐽′+1)𝜆3× [𝜒𝑒𝑑𝑆𝑒𝑑(𝐽′: 𝐽) + 𝜒𝑚𝑑𝑆𝑚𝑑(𝐽′: 𝐽)] (2.19)
If there is more than one transition related with level Jʹ, the ratio among each
transition can be determined by:
𝛽(𝐽; 𝐽′) =𝐴(𝐽;𝐽′)
∑ 𝐴(𝐽;𝐽′)𝐽′ (2.20)
and the radiative transitions rates can be calculated by:
𝑊𝐽𝑅 =
1
𝜏𝑅= ∑ 𝐴(𝐽; 𝐽′)𝐽′ (2.21)
Although the unambiguous interpretation of J-O parameters has always been
challenging, in general, it is agreed that smaller values of the Ω2 parameter indicates
chemical bonds that have ionic rather than covalent character, while low Ω4 and Ω6
values may be due to high rigidity of the host matrix [121]. Also it is believed that
35
parameter Ω2 is related with asymmetry of the crystal field on the rare-earth ion, the
larger the Ω2 the more asymmetric the glass host [169].
Typical values for J-O parameters in some glasses are summarized [169] in Table
2.3. (Added information a is cited from ref [170], b is from [136], and c is from [171]).
Table 2.3 Typical value of Judd-Ofelt parameters in some glasses.
Glass Ω2 (x10-20 cm2) Ω4 (x10-20 cm2) Ω6 (x10-20 cm2)
Fluorophosphate 2.90 1.63 1.26
Germanate 4.81 1.41 0.48
Silicate 4.23 1.04 0.61
Aluminate 5.60 1.60 0.61
Fluoride 2.90 1.27 1.11
Phosphate 3.89 1.01 0.55
Tellurite 5.05 1.45 1.22
Bismuth 3.86 1.52 1.17
Chalcogenide (Ge-Ga-S)a 6.9±0.1 2.19±0.02 1.05±0.01
Chalcogenide (Ge-As-S)b 6.98 2.53 0.78
Chalcogenide (GeGaSbS)c 8.7 2.5 1.4
Amongst all the rare-earth ion host materials listed in Table 2.3, the big Ω2 values
chalcogenide glasses have imply their covalent character of chemical bonds. Also, the
relatively high Ω4 values of chalcogenide glasses indicate the less rigidity of the host
matrix. As a result of the widely varying J-O values for the same rare-earth ion in
different host materials, the emission parameters also vary widely. Taking Er3+ for
example, in silicate hosts the 4I13/2 energy level lifetime can exceed 10 ms [172], while
in chalcogenide glass host the number drops to around 4 ms [119], but still long enough
for applications such as amplifiers and lasers. In Table 2.4, the reported emission
properties of Er3+ in different glass hosts are summarised.
36
Table 2.4 Emission properties of 4I13/2→4I15/2 transition of erbium ion in different host
materials [2, 154, 173], and FWHM refers to full width at half maximum.
Glass hosts Refractive
Index at
1.5 µm
Peak
emission
cross-
section /
cm2
Emission
FWHM
bandwidth
/ nm
Lifetime
/ ms
Maximum Er3+
concentration
Max
phonon
energy
/ cm-1
Silica ~1.5 7x10-21 20 12 ~3x1019 ions/cm3 1200
Alumina ~1.7 6x10-21 55 [174] 8 >3.4x1020
ions/cm3 [175]
1000
Aluminosilicate ~1.6 6x10-21 43 10 500 ppm
(~1x1019
ions/cm3)
1100
Phosphosilicate ~1.6 6x10-21 27 10 2.5 at% (~5x1020
ions/cm3)
1200
Tellurite [176,
177]
~2.1 13x10-21 80 4 1x1021 ions/cm3
(10 mol% in
TeO2-WO3-
La2O3 [178] )
700
Chalcogenide
(sulfide based)
[130]
~2.4-3 15x10-21 45 2.5 1x1019 ions/cm3 400
ZBLAN [179] 1.5 4.2x10-21 82 9 7.9x1020
ions/cm3
500
Among these hosts, erbium ion doped chalcogenide glass has the highest refractive
index, emission cross-section, and the lowest maximum phonon energy, along with
reasonable full width at half maximum (FWHM) bandwidth and erbium ion solubility.
All of these make chalcogenide glasses competitive candidate for erbium ion doped
amplifier/lasers.
37
2.3 Energy migration processes in rare-earth ion doped
materials
2.3.1 Ion-Ion interactions
At low dopant concentration, the average distance between two adjacent ions is
sufficiently large that the interactions between separate ions can be neglected. As the
dopant concentration increases, the average inter-ion distance decreases, and thus ion-
ion interactions involving excited state energy transfer start to occur. Such interactions
can occur between rare-earth ions of the same or different species. Within the same
rare-earth ion type, ion-ion interaction leads to either a non-radiative loss mechanisms
or luminescence from unwanted transitions. Between different rare-earth ions, it can
offer new pumping schemes to extend the selection of pumping sources, or a means to
de-excite targeted energy levels in four or more state lasing mechanisms [154], and also
in 3 level systems, e.g. Ce3+ & Eu3+ with Er3+ in TeO2 for 980 nm pumping [180, 181].
Ion-ion interactions due to multipolar interactions between neighbouring rare-earth
ions have been carefully studied. Based on Auzel’s research [149], the probability for
energy transfer between two ions at a sufficiently large distance R can be written as:
𝑃 =1
𝜏𝑠(𝑅0
𝑅)6
(2.22)
where τs is the actual lifetime of the sensitizer (the ion which is being first directly
excited) excited state, including multiphonon non-radiative decay, and R0 is the critical
transfer distance for which excitation transfer and spontaneous deactivation of the
sensitizer have equal probability.
However, Dexter pointed out that this theory could be extended to include higher
multipole and exchange interactions [182]. Thus, the energy transfer probability for
electric multipolar interactions can be modified as follows:
𝑃 =1
𝜏𝑠(𝑅0
𝑅)𝑠
(2.23)
where s=6 for dipole-dipole interactions, s=8 for dipole-quadrupole interactions, and
s=10 for quadrupole-quadrupole interactions [149].
Normally, in a single rare-earth ion containing system, a number of different ion-ion
interactions may occur, the most important of which are outlined in the following
38
section. The erbium ion is chosen as an example to illustrate the mechanism of each
process.
2.3.2 Up-conversion process
Up-conversion emission is an important phenomenon in rare-earth ion doped materials.
Emissions of higher energy than that of the excitation light are seen, which is due to the
so called anti-Stokes emission or up-conversion process. There is quite a number of
possible up-conversion mechanisms in rare-earth ion doped materials, of which the most
important include: energy transfer up-conversion (ETU), excited state absorption
(ESA), cooperative sensitization, cooperative luminescence and 2-photon absorption
excitation [149, 154].
(a) Energy transfer up-conversion (ETU)
The process of up-conversion by sequential energy transfers, also named APTE (for
‘addition de photon par transferts d’energie’) by Auzel [149]. In this case, energy
transfers from one ion already in some excited state to another excited ion, which may
be in a different excited state, to promote it to a higher energy state. Resonance or
phonons may be involved to help match the energy difference in a real case. In an Er3+
ion system for example, with 1490 nm excitation, two ions are excited into the 4I13/2
metastable state. If ETU occurs, then one ion will transfer its energy to the other and
decay to the ground state, whilst the one receiving the energy will be promoted to the
higher energy level 4I11/2, corresponding then to the 800 nm emission. The process can
also apply to ions in different excited states, for example, an ion in the 4I11/2 level can
receive energy from an excited ion in the 4I13/2 level and be further promoted to the 4S3/2
level to radiate 530 nm. The schematic energy level diagram for ETU process is shown
in Figure 2.2 (a).
Figure 2.2 Schematic energy level diagram showing the ETU (a) and ESA (b) [149].
39
In ETU then, only the matching of the energy difference between the upconversion
level and the de-excitation levels involved matters rather than the absolute energy,
therefore this process applies potentially to many up-conversion wavelengths [149].
(b) Excited state absorption (ESA)
In this process, an ion in an excited state absorbs another pump or signal photon, and is
further excited into a higher energy level. Obviously, this process requires the
availability of an energy level which is roughly a pump or signal photon energy above
the current level to occur. For example, in the Er3+ system, an ion in the 4I13/2 state
absorbs another 1490 nm photon and promotes to the 4I9/2 state. The upper state excited
ion could return to the metastable level by multiphonon relaxation or radiative decay, or
back to the ground state with a high energy photon emission, 800 nm emission in the
case of Er3+. The schematic energy level diagram for this process could be found in
Figure 2.2 (b). ESA reduces the excitation efficiency as pump photons are lost to either
heat or emission at unwanted wavelengths. In the case of signal photons, it can also
clamp the extent of the gain window. Examples of this process are shown in ref. [183,
184].
(c) Cooperative sensitization
In cooperative sensitization, two ions in the metastable state give their energy to a third
ion thereby decaying to the ground state. The third ion which accepts the energy will be
promoted to an excited state around the sum of the energies of the de-excited ions from
its ground state. An Yb3+/Tb3+ co-doped system makes a good example of this process.
Tb3+ ions at ground state accept energy from two Yb3+ ions of the 2F5/2 excited states,
and then the Tb3+ ions is promoted to the 5D4 excited state, while the two Yb3+ ions
return to their 2F7/2 ground state. Example of this process could be found in ref [185].
Figure 2.3 Schematic energy level diagram showing the cooperative sensitization (a),
cooperative luminescence (b) and 2-photon absorption excitation (c) [149, 154].
40
Cooperative luminescence is a radiative process in which two metastable level ions
simultaneously de-excite to emit one photon with the sum of their excitation energies
[149]. Two-photon absorption excitation is that where the ion absorbs two photons
simultaneously to promote it into a higher energy state. Due to the lack of interaction
with the medium, the efficiency of these two processes is relatively low. Research on
these phenomena can be found in ref [186]. Schematic energy level diagrams of
cooperative sensitization, cooperative luminescence and 2-photon absorption excitation
are shown in Figure 2.3.
From Auzel’s reports [149], the relative efficiency of ETU is much higher than the
other processes, almost two orders of magnitudes higher than the efficiency of two-step
absorption and three orders of magnitudes higher than that of cooperative sensitization.
The efficiency of different up-conversion processes in some materials is listed in Table
2.5.
Table 2.5 Comparison of efficiency of different 2-photons up-conversion processes [187]
Mechanism (section) Efficiency /cm-2 W-1 Example
A.P.T.E (ETU) [2.2.2(a)] ~10-3 YF3:Yb3+:Er3+
Excited state absorption (ESA) [2.2.2(b)] ~10-5 SrF2:Er3+
Cooperative sensitization [2.2.2(c)] ~10-5 YF3:Yb3+:Tb3+
Cooperative luminescence [2.2.2(c)] ~10-8 Yb3+ PO4
Second-harmonic generation (SHG) ~10-11 KDP
2-photon absorption and emission [2.2.2(c)] ~10-13 CaF2:Eu2+
KDP refers to Potassium Dihydrogen Phosphate (KH2PO4).
Although each up-conversion process has its own mechanisms, in reality, different
up-conversion processes may occur simultaneously, and the final emission is always a
combination of all processes involved.
2.3.3 Energy transfer
An ion in the metastable state may also interact with a nearby ground state ion,
promoting it to the metastable state by passing over its energy. This can occur with both
radiative and non-radiative pathways. In the radiative pathway, one ion decays from the
41
metastable state to the ground state by emitting a photon which is immediately absorbed
by the adjacent ground state ion promoting it into the metastable state. Schematic
energy level diagram for this process is shown in Figure 2.4(a). This can cause a
distortion of the emission spectrum as energy can be lost or gained [154] and also the so
called ‘radiative trapping’ phenomenon that leads to a longer apparent lifetime
especially in highly doped or large-volume samples [188]. On most occasions, the non-
radiative energy transfer mechanism is more important, in which the energy transfers
without emission and absorption of photons. Phonons are often involved to bridge the
energy gap when the energy between related ions is not equal. This interaction can be
explained by either the short-range exchange or long-range electric multipolar
mechanisms [188]. Although radiative emission may still occur from the second ion that
receives the energy, the probability of non-radiative decay is increased with each
successive transfer, and hence this is always considered as a loss mechanism.
Figure 2.4 Schematic energy level diagram for energy transfer (a), and cross-relaxation (b),
in Er3+ ion [154].
2.3.4 Cross relaxation
Cross-relaxation is the process where an ion at a higher excited state gives part of its
energy to another ion at the ground state such that both end up in a metastable state.
Schematic energy level diagram for this process is shown in Figure 2.4(b). Taking
erbium ion as an example, the energy gap between the 4I9/2 state and the 4I13/2 state is
similar to the gap between the 4I13/2 state and the 4I15/2 state, thus an ion in the 4I9/2 state
may give part of its energy to an ion in the 4I15/2 ground state and then decay to the 4I13/2
state, whilst the ground state ion receives the energy and is promoted to the 4I13/2
metastable state. Obviously, this process may increase the population of the metastable
42
state by the decay of ions from the higher excited state and also ion promotion from the
ground state.
2.4 Concentration quenching
Concentration quenching is usually the term used to sum up the effects of the previously
described ion-ion interactions which result in the decrease of rare-earth ion fluorescence
efficiency with increasing concentration of the ion. In highly doped samples, the
average distance between neighboring ions is short, thus the ion-ion interaction
processes mentioned above (see Section 2.3) occur and the efficiency of fluorescence is
degraded. Concentration quenching also manifests itself as a shortening of the
metastable state lifetime, and the relation between lifetime and concentration has been
expressed by an empirical formula [189, 190]:
𝜏𝑜𝑏𝑠 =𝜏0
1+(𝜌/𝑄)𝑝 (2.24)
where τobs is the observed fluorescence lifetime, τ0 is the ideal fluorescence lifetime with
zero concentration, ρ is the rare-earth ion concentration, Q is the quenching
concentration and p is a phenomenological parameter characterizing the steepness of the
corresponding quenching curve. Normally, the concentration Q with a lifetime τ equal
to τ0/2 is defined as the quenching concentration in this material. Quenching
concentrations of Er3+ in several different hosts are listed in the Table 2.6 [189].
43
Table 2.6 Quenching concentrations of Er3+ in several different hosts.
Type of glass Q / ions cm-3 τ0 / ms p
Sol-gel silicate 80SiO2-20TiO2-0.5Yb2O3-10Al2O3
waveguide (mol%) [189]
1.4x1020 6.4 1.27
Al2O3-GeO2-SiO2 fibre [191] 1x1020 9.6 1.27
Borosilicate [192] 8x1019 9.9 1.55
Phosphate [193] 2-3x1019 14 2
Aluminate 58.5GaO-27.5Al2O3-8.4MgO-5.6SiO2 (wt%)
[194]
4x1018 7.6 0.97
Germanate 70GeO2-24PbO-6PbF2 (wt%) [194] 5.2x1018 7.3 1.74
Sulfate Ga2S3-GeS2-La2S3 [195] 3.2x1020 2.2 1.29
Alkali-borosilicate BGG 31 [195] <5.8x1019 >0.96 unknown
Ge25Ga10Se65 (at%) (data from this work) 3.6x1020 1.69 1.415
In this table BGG 31 refers to BaO-GeO2-Ga2O3 glass, Q is the quenching concentration, τ0 is the ideal
fluorescence lifetime with zero concentration and p is the phenomenological parameter characterizing the
steepness of the corresponding quenching curve.
Amongst the host materials, chalcogenide glasses have the shortest lifetime due to
their high value of Judd-Ofelt parameters and high refractive index [196]. In Ga
containing chalcogenide glasses the quenching concentration is one of the highest,
which means high dose rare-earth ions can be doped in Ga containing chalcogenide
hosts to achieve high gain in a short length, which is favourable for compact devices.
2.5 Solubility of rare-earth ion in glasses
Rare-earth ions in any host may form crystalline or amorphous precipitates when the
concentration exceeds a critical value. This can be in the form of clusters of rare-earth
ions in compounds, alloys or glasses within the host matrix. Such processes are
detrimental to photoluminescence either by ion-ion interactions between essentially
bonded rare-earth ions or the formation of other compounds that are not optically active
in the host. Rare-earth ion solubility varies with host material. Maximum concentrations
for silicate glasses and phosphate glasses are in the range of 4-9x1020 ions/cm3 [154].
44
Tellurite glass is known to have high rare-earth ion solubility and can contain erbium
ion of around 5 at% (~1x1021 ions/cm3) without any evidence of clustering [176]. As the
workhorse of the chalcogenide glass system, bulk As2S3 is considered to have a low
erbium ion solubility [197]. In Gu et al.’s report, bulk samples of Er3+ doped As2S3 with
nominal Er3+ concentrations from 0.02 to 4 wt% (~6x1020 ions/cm3) were prepared.
Unfortunately, in the X-ray pattern, the sample with the highest Er2S3 dopant
concentrations exhibited weak, sharp-line X-ray diffraction pattern which were
consistent with the polycrystalline Er2S3 source materials, indicating the existence of
precipitates and cluster in the host material [143]. Amongst the family of chalcogenides,
gallium-containing materials have been proven to have better rare-earth ion solubility.
In the Ge-Ga-Se/S system, Er3+ ion of up to 2 at% (~ 2x1020 ions/cm3) illustrated that
erbium ion could be well dispersed and optically active [198, 199], implying the
potential in compact planar waveguide amplifiers.
2.6 Conclusion
This chapter covered the basic theory of rare-earth ion doped materials. Concepts
related to rare-earth ion doped materials such as lifetime, absorption/emission cross-
section and maximum pump efficiency were discussed in detail. Judd-Ofelt theory was
introduced as a tool to predict the transition probabilities or oscillator strength of
transitions between energy levels. Energy migration processes in rare-earth ion doped
materials were catalogued by their mechanisms. Concentration quenching and solubility
of rare-earth ion in glasses, including chalcogenide glasses, were discussed.
45
Chapter 3
Co-thermal evaporation of erbium ion doped
As2S3 waveguides
3.1 Diarsenic trisulfide: As2S3
3.1.1 Background of As2S3
Diarsenic trisulfide (also known as and hereafter referred to as arsenic trisulfide), As2S3,
as a member of chalcogenide family, has been well studied in the past several decades.
Besides the generic chalcogenide properties it possesses, as a bulk glass it has
particularly good chemical stability and resistance to crystallization, a high optical
damage threshold, usefully large optical nonlinearity, a relatively large optical bandgap,
and can easily be deposited by thermal evaporation. These characteristics have made it
the subject of many studies since the 1970s in bulk glass, fibre and waveguide forms.
The first study resulted from the desire to make glass with transparency beyond 3-5
µm, initially using sulfur-based glasses by Frerichs in the 1950s [200]. Jerger and Fraser
of Servo Corporation then developed useful As-S based glasses in 1953 [201].
The first As-S core-clad fibres were fabricated by Kapany and Simms in the early
1960s, though with a relatively high loss of around 20,000 dB/km at 5.5 μm [202]. The
loss of chalcogenide fibres was subsequently reduced significantly over the intervening
years to 35 dB/km at 2.44 µm in an As2S3 unclad fibre in Kanamori et al.’s report of
1984 [203].
As2S3 also became the workhorse material for studies into planar waveguide
nonlinear optics e.g. [4, 62, 64, 66, 204-206]. Various methods have been used to realise
high quality waveguides based on As2S3 films. In Zoubir et al.’s report, channel
waveguides were laser written in As2S3 thin films [207]. By using focused femtosecond
pulses from a 25-MHz repetition rate Ti:sapphire laser, 2 cm long waveguides were
fabricated with refractive index differentials Δn>10-2. The obtained waveguides showed
reasonable light confinement ability by confining more than 70% of the launched light
exiting the film end face in the channel waveguide. Wet etching is another way to
fabricate chalcogenide glass waveguides. In Su et al.’s report, a focused laser beam was
applied to form waveguide patterns on an As2S3 film deposited using thermal
46
evaporation [208]. The exposed and unexposed As2S3 thin films had different etching
rates in I2 and CS2 solutions, as a result, a ridge waveguide structure was obtained with
this method. However, rough surfaces were observed on the scanning electron
microscope (SEM) image of the etched ridge waveguide, leading to an average loss of
5.5±0.5 dB/cm at 1330 nm. As2S3 waveguides fabricated through hot embossing were
demonstrated in [209]. The propagation loss of such waveguides was reported at 0.52
dB/cm for the transverse electric (TE) polarization and 0.41 dB/cm for the transverse
magnetic (TM) polarization at 1550 nm for a waveguide cross-section dimension of 3.8
× 1 μm. Dry etching has proven to be a promising way to fabricate chalcogenide
waveguides. Ruan et al. reported As2S3 rib waveguide losses as low as 0.25 dB/cm at
1550 nm achieved using dry etching with CF4 and O2. With the obtained waveguides, a
phase shift ~π due to self phase modulation has been obtained using a ~40 W peak
power pulse in a 5 cm long waveguide [12]. Madden et al. using a different dry etching
recipe (using CHF3 instead of CF4), and obtained As2S3 rib waveguides with loss <0.05
dB/cm, which is the current lowest loss in chalcogenide waveguides [13].
3.1.2 Properties of As2S3
(a) Optical transparency
The transmission spectrum of a 37 mm optical path length As2S3 sample and of a 2 mm
optical path length As2S3 plate are shown in Figure. 3.1. The useful transmission
window for planar devices made from As2S3 is from ~0.8-8 µm. At the short
wavelength side, below 0.7 µm, the Urbach absorption edge and the weak absorption
tail (WAT) occur inducing a fast increase in absorption. Also an absorption band at 4
µm, corresponding to the S-H vibrations, appears in the curve as well.
47
Figure 3.1 Absorption edge from a 37 mm optical path length As2S3 sample (a);
transmission spectrum (uncorrected for reflection losses) from a 2 mm optical path length
As2S3 sample (b) [210].
(b) Refractive index
The refractive index of bulk As2S3 is 2.438 at 1.53 µm, and drops with increase of
wavelength [211]. The high refractive index enables strong optical field confinement
and hence small waveguide bend radius and small optical mode area with enhanced
optical intensities. These are critical characteristics for application in integrated optics
and efficient nonlinear interactions. The zero dispersion wavelength for As2S3 material
lies in the MIR at 4.9 µm, and there is strong normal dispersion at telecommunication
wavelengths (1.55 µm) of ~-360 ps nm-1 km-1. Thus dispersion engineering is normally
required to shift the zero dispersion wavelengths in practice for telecoms NLO
applications [212]. Figure 3.2 shows the refractive index and dispersion curve for As2S3.
Figure 3.2 Refractive index and dispersion curve of As2S3 [59].
48
(c) Raman spectrum
Raman spectroscopy provides information about bond structure via molecular
vibrations, and thus is widely used for sample identification and quantification in
materials science. A typical Raman spectrum of bulk As2S3 glass has been reported in
the literature [213-216], and Raman spectrum of an in-house made bulk As2S3 glass
under 830 nm excitation measured in our lab is shown in Figure 3.3. The spectrum is
strongly dominated by the band at 345 cm-1 attributed to the symmetric stretching
vibrational mode of AsS3/2 pyramids. Besides this strong bond, there are two weak
shoulders at 315 cm-1 and 390 cm-1 on both sides of the 345 cm-1 band, which can be
associated with asymmetric stretching modes of AsS3/2 pyramids and As-S-As bridges,
respectively. In the small Raman shift range, small bands at 190 cm-1 and 235 cm-1
corresponding to weak bonds that can be attributed to the bending modes of AsS3/2
pyramids and S8 and As4S4 molecules. The small band around 495 cm-1 associated with
the S-S stretching vibration in S8 rings indicates the presence of excess sulfur in this
bulk glass.
Figure 3.3 Raman spectrum of an in-house made bulk As2S3 glass under 830 nm excitation.
3.2 Erbium ion doped As2S3
3.2.1 Research in erbium ion doped As2S3 glasses
As the workhorse material for planar waveguides in the chalcogenide glass family, rare-
earth ion doped As2S3 for amplification and lasing applications in the NIR and MIR
range was promising and studied for decades. Besides properties that are preferable for
a rare-earth ion dopant in a chalcogenide glass, such as low phonon energy (350 cm-1)
49
and wide transparency window, As2S3 has the advantage of a band gap in the green,
meaning all tail absorption processes are low at common rare-earth ion pump
wavelengths (e.g. 808, 980, 1064 nm). Unfortunately, bulk As2S3 glass cannot be doped
without clustering at high enough concentrations for planar waveguide amplifiers [118,
197] to date blunting the advantages offered by the material. Thin films, however, offer
an alternative perspective as they are formed under highly non-equilibrium conditions
where there is insufficient time and thermal energy for clusters to form in the deposited
film, leading to new opportunities for doped devices.
There are investigations on rare-earth ion (and particularly erbium ion) doped As2S3
films and waveguides [128, 130, 217]. Fick et al. investigated Er3+ doped As2S3 films
formed using thermal evaporation and subsequent ion implantation [130]. Emission
cross-sections up to 1.6x10-20 cm2 were measured for the 4I13/2→4I15/2 transition that is
two times higher than for Er3+ doped silica glass. A 2.3 ms lifetime of the 4I13/2
metastable state under 983 nm pumping was achieved from planar samples with Er3+
concentration of ~5x1018 ions/cm3. Due to the low phonon energies of chalcogenide
glasses, a relatively long lifetime of the 4I11/2 excited state was measured of the order of
0.25 ms. Improvements in PL intensity of up to almost 10x were seen in some cases
after annealing the films just below the glass transition temperature at 165 °C for 2
hours. Lifetime was also improved by the post implant anneal but only by about 10%.
Highly erbium ion doped films (~4 at%) were produced using co-evaporation of
As2S3 and Er2S3 by Lyubin et al. [128]. Emission at 1.54 μm, was about 3x the intensity
that of similarly doped GeGaS bulk glass. A linear PL intensity dependence on the
excitation laser intensity was observed. Photostructural transformation phenomena were
also studied by using an installation with two linearly polarized lasers, an.Ar+ laser (488
nm, 10 mW), acting as the inducing beam, and a He–Ne laser (633 nm, 0.1 mW)
working as a probing beam. They concluded that although the transparency of the Er3+-
doped film decreased, the photodarkening effect in the doped film was smaller than in
the non-doped film.
Fuchs et al. [217] studied the spectral properties of Er3+ doped As2S3 films deposited
using RF sputtering. An As2S3 glass disc with a piece of erbium metal partly covering
its surface was used as sputtering target. After thermal annealing at 150 °C, a 4 ms PL
decay lifetime of the 1.55 μm emission was measured when pumped at 977 nm,. The
annealing also led to an increase in the PL intensity of up to 40x when measured in slab
waveguide geometry. Erbium ion concentrations were estimated at ~1 at% for the
50
highest doped sample which had inferior PL performance compared to the lower doped
samples with unknown concentrations.
Despite these promising results, and other work on alternative chalcogenide host
materials, there has been, however, no report of amplification in erbium ion doped
chalcogenide glass devices. To realise amplification in chalcogenide glasses, the
emission properties in rare-earth ion doped chalcogenide glass must be well understood
and available methods used to fabricate low loss waveguides. The Laser Physics Centre
at the Australian National University has significant expertise in depositing and
processing As2S3. As stated above (see Section 3.1.1), an As2S3 rib waveguide with
record low loss of 0.05 dB/cm was achieved by photolithography and plasma etching
[13] providing a good base for rare-earth ion doped amplifier studies in As2S3. Given
the facilities and experience available, it was decided to make the first step with erbium
ion doped As2S3.
3.2.2 Fabrication of bulk As2S3 glass
The As2S3 used here was fabricated using the conventional melt-quenching method,
either in house or commercially sourced from Amorphous Materials Inc. In this method,
high purity arsenic and sulfur (5N) were weighed inside a dry nitrogen glove box and
loaded in the required proportions into a pre-cleaned quartz ampoule. The ampoule was
then sealed under vacuum (~10-4 Pa) using an oxygen-hydrogen torch, and introduced
into a rocking furnace for melting of the contents at ~800 °C. The melt was then
homogenized for periods typically exceeding 12 hours, before the ampoule was
removed from the rocking furnace and water-quenched. The resulting glass boule was
subsequently annealed at a temperature typically 20 °C below its glass transition
temperature (Tg) of ~180 °C, before being slowly cooled down to room temperature. A
piece of bulk As2S3 made in-house (without elemental purification) is shown in Figure
3.4. The starting elements may also be purified before glass making to remove oxygen,
water, and carbon contaminants. Details on starting element purification can be found in
ref [218].
51
Figure 3.4 Typical in-house made bulk As2S3 glass.
3.2.3 Erbium ion doped As2S3 film deposited using co-thermal evaporation
Different methods, such as RF sputtering [219], pulsed laser deposition [220, 221],
chemical vapour deposition (CVD) [222], solvent based spinning [223, 224] or thermal
evaporation [225] may be used for film deposition. Each method has its own advantages
and drawbacks. Among these approaches, thermal evaporation is one of the most
popular film fabrication methods for low loss waveguide devices [4, 13]. However, for
most ternary and quaternary glasses (and especially those containing gallium and rare-
earth ions), a large contrast in the saturated vapour pressure of these elements and their
various compounds in the melt lead to severe difficulty in controlling the film
composition using thermal evaporation. This usually leads to significant differences
between the properties of as-deposited films and the bulk glass, often producing nearly
rare-earth- [128, 130], and Ga-free films. Besides the stoichiometry of the film, physical
parameters like refractive index of as-deposited films were found to be markedly
different from bulk counterparts [225]. This also results in the films being highly
photosensitive [213]. The significant difference between as-deposited films and bulk
glass is likely because films are fabricated in non-equilibrium conditions. Therefore, an
approach that can release the excess internal energy stored in the films and relax them
back to the thermal metastable equilibrium state is required.
It has been reported that at room temperature the relaxation time for non-annealed
As2S3 films made by electron beam evaporation to the equilibrium state is of the order
of one year [226], but thermal annealing can accelerate this procedure. Also thermal
annealing can remove the photo-structurally-induced photosensitivity, but residual
photosensitivity remains, the origin of which is still not entirely clear [227, 228]. Also
one should note that so far no film with properties identical to its bulk counterpart has
been fabricated [229] [213].
52
As2S3 bulk glass cannot be doped sufficiently with rare-earth ion to avoid clustering,
so whilst it evaporates above Tg [13, 227, 230], it is not a suitable starting point for
erbium doped thin films. As a result co-thermal evaporation method was instead chosen
to dope the film. Crushed As2S3 glass and an erbium source (Er metal pieces in this
Chapter, Er2S3 and ErCl3 in Chapter 5) were used in film deposition. The two sources
had their own shutters and evaporation rate monitors, and thus theoretically any desired
doping concentration could be achieved. An internal schematic of the co-thermal
evaporation chamber manufactured by Angstrom Sciences is shown in Figure 3.5. Here
the resistively heated boat at the bottom centre of the diagram was used to evaporate the
As2S3, and one of the smaller furnaces in the group of five in the centre of the diagram
was used for the erbium.
Figure 3.5 Internal view of the co-thermal evaporation chamber manufactured by
Angstrom Sciences: 6 sources for co-thermal evaporation, and 3 guns for RF sputtering.
In the first experiment, erbium metal was used as the dopant. High purity As2S3
(made from 5N elements) and erbium metal were loaded into resistively heated sources.
Each source had one quartz microbalance thickness monitor overhead to monitor the
evaporation rate, and another monitor was placed above the shutter at the same height as
the wafer to check the film thickness. Evaporation was performed at a vacuum level
around ~1x10-5 Pa and with a source-wafer distance of ~40 cm. After reaching a
pressure set point of ~6x10-4 Pa, the temperature of each source was slowly increased to
~280 °C over a 20 minute period to bake out the evaporants before ramping to the
53
evaporation temperature and setting the individual rates. The erbium metal source was
also ‘cleaned’ by heating up at evaporated temperature for 10-15 minutes with shutter
closed to remove the possible contamination due to surface oxidation. Typical
evaporation rates for As2S3 and erbium were fixed at 0.1 nm/s and 0.001 nm/s,
respectively. After the desired evaporation rates were achieved, the shutter covering the
wafers was opened and the evaporation started. When the film thickness read from the
film thickness monitor reached the desired thickness set point, the system automatically
closed all the shutters and ramped down the power for both sources gradually and
allowed them to cool to <100 °C before venting the chamber.
It is well known that as-deposited As2S3 films degenerate easily through
crystallization via surface diffusion under the influence of room light [231, 232],
therefore, as soon as the films were removed from the chamber, a ~100 nm thick SU-8
(SU-8 2 from MicroChem Corp. diluted 1:4 with SU-8 GBL developer) layer was spin-
coated and ultraviolet-cured (UV) on the film to prevent As2S3 evaporation at high
temperature during the following thermal post-treatment and As2O3 crystallite formation
on the surface [233, 234].
As previously noted, one of the issues in deposited films is that the as-deposited
films are not identical with their bulk glass counterparts in terms of refractive index,
bandgap, etc. [213, 225]. As a result of the highly non-equilibrium deposition
conditions, thermally-evaporated As2S3 films tend to contain significant chemical and
structural disorder such as As–As and S–S homopolar bonds, which are at lower
concentrations in stoichiometric bulk glasses [213, 235].
Thermal annealing is a simple way to accelerate the relaxation of a film towards its
metastable equilibrium state. During heat treatment, it has been reported that the
properties of films tend to change and finally reach a condition close to their bulk
counterpart [213]. In this research, thermal treatment of as-deposited films was carried
out at temperatures of 130 °C for 24 hours in a vacuum oven (~10 Pa), which is below
the glass transition temperature (Tg ~180 °C) and low enough to prevent cracking of the
film after thermal treatment due to the large mismatch in thermal expansion coefficient
of As2S3 (~23x10-6/°C) and SiO2 (~0.5x10-6/°C).
Besides thermal treatment, it was reported that optical treatment with light
illumination around the optical bandgap of the films, evolves the structure of the film in
a similar way to thermal treatment [234]. Light treatment was carried out in the
following way: the emission of an array of 100 W tungsten halogen reflector spotlamps
54
was filtered using green band-pass gel filter (G28 from Lee Filters) to produce a
spectrum extending from 500-570 nm which overlapped with the bandgap of As2S3
films. Then full 4-inch diameter wafers were exposed under the uniform light with an
illumination intensity of 2-3 mW/cm2 for 48 hours in the ambient atmosphere. In later
experiments, the halogen lamps were replaced with green LED floodlights providing a
uniform illumination ~8-10 mW/cm2.
The thickness and linear refractive index (n) of the films were measured using a
spectroscopic reflectometer (SCI FilmTek 4000) applying a Tauc-Lorentz model to fit
the acquired data. In the FilmTek 4000 system, spectroscopic reflection data (400-1650
nm) were gathered at two different angles (normal incidence and ~70°incidence). The
Lorentz oscillator parameters and the Tauc bandgap were the free variables used in a
nonlinear fitting procedure to least squares fit to the measured reflectance data, and an
iterative error minimization procedure employed to fine tune the model parameters and
thickness. The system was capable of 1 part in 105 index/thickness accuracy for a layer
of silica on silicon. Therefore, with this system, important film parameters such as
thickness, optical bandgap of the material, and refractive index and dispersion from
400-1650 nm could be obtained.
The as-deposited, thermal and light treated films were tested in terms of thickness,
refractive index at 1550 nm and bandgap, using the SCI FilmTek 4000 and the results
are summarized in Table 3.1.
Table 3.1 Properties of As2S3 films with different treatment methods.
Type of treatments Thickness /
nm
Refractive index
@1550 nm
Bandgap /
eV
As-deposited 892 2.35 2.325
Thermal treatment (130 °C for 24 hours) 877 2.41 2.337
Light treatment (green light with
intensity of 2-3 mW/cm2 for 48 hours) 881 2.43 2.265
From the results, the film thickness decreased from 892 nm for an as-deposited film
down to 877 nm and 881 nm after thermal or light treatment, respectively. This change
is believed to be due to the consolidation of the films during the thermal or light
treatment. The refractive index at 1.55 µm increased from 2.35 for an as-deposited film
to 2.41 in the thermally treated sample, and 2.43 for the post-light-treatment sample,
which is almost the refractive index of their bulk counterpart, indicating that both
55
thermal and light treatment offered the similar effects that brought the as-deposited film
back towards the bulk condition. As for the energy of bandgap, thermal treatment gave
film a tiny increment, while the light treatment dropped it from 2.325 eV down to 2.265
eV, the reasons for this are unclear at present.
3.3 Characterization of erbium ion doped As2S3 films
3.3.1 Raman spectrum of erbium ion doped As2S3 film deposited using thermal
evaporation
For Er3+ doped As2S3 films, the Raman spectrum was measured to check whether there
were any structural changes between the doped film and its un-doped and bulk
counterparts. The Raman spectrum of a 600 nm thick film with 1.8x1020 ions/cm3 Er3+
doping (thermally treated at 130 °C for 24 hours) on a 1.5 μm thick thermally oxidized
silicon (TOX) substrate was measured in a Horiba Jobin Yvon 64000 spectrometer
system with a 50x NIR objective with numerical aperture (NA) of 0.75. A 630 nm laser
was employed as the excitation source, and a charge-coupled device (CCD) detector
installed on the spectrometer recorded Raman spectra. The power of the 630 nm
excitation laser was set ~1 mW to avoid damage the sample surface. For comparison,
bulk As2S3 glass and an As2S3 film without erbium ion doping (thermal treated at 130
°C for 24 hours) were also measured and the spectra are shown in Figure 3.6 with
vertical offsets to aid viewing.
Figure 3.6 Raman spectra of: bulk As2S3 glass; As2S3 film (thermal treated at 130 °C for 2
hours) and Er3+:As2S3 film (thermal treated at 130 °C for 2 hours), the spectra were
recorded under 630 nm excitation.
Due to the small thickness of film, the signal to noise ratio dropped and a much
longer acquisition time was required to get a reasonable spectrum. The band at 345 cm-1,
56
attributed to the symmetric stretching vibrational mode of AsS3/2 pyramids, still
dominates the whole spectrum. The shoulder at 315 cm-1 is hard to distinguish, while
another shoulder around 390 cm-1 associated with asymmetric stretching modes of
AsS3/2 pyramids and As-S-As bridges is still traceable. Weak bands at 190 cm-1 and 235
cm-1 corresponding the bending modes of AsS3/2 pyramids and S8 and As4S4 molecules
are still clear to see, but, it is obvious that the band at 235 cm-1 is stronger in films than
in the bulk glass counterpart, indicating more S8 and As4S4 structure are formed during
film deposition. This is common in thin films [236, 237]. The Raman spectra from films
with and without 0.6 at% erbium ion are very similar within the limits of low signal to
noise ratio in Figure 3.6, implying the added erbium ions have no significant effect on
the molecular structure of film.
3.3.2 Optical properties of erbium ion doped As2S3 film
(a) Lifetime and PL spectrum of erbium ion doped As2S3 films
There are two types of lifetime that are used to describe the decay curve. One is the 1/e
lifetime of the captured decay curve, which is defined as the time to 36.8% of the
fluorescence intensity after the pumping is turned off and includes all effects impacting
the radiative decay. The other is the underlying (hereafter referred to as intrinsic)
radiative lifetime, which is the lifetime the fluorescence has in the absence of population
dependent processes, though this is still affected by some ion-ion interactions it is the
best measurable approximation to the actual radiative lifetime.
To accurately measure the lifetime of thin films, an all-fibre confocal setup first
proposed in [152] was used, the scheme of which is shown in Figure 3.7. Pumping
power at 1490 nm or 980 nm was delivered to the edge of the film via a lens tipped fibre
with 14 μm working distance. The lensed fibre produced a 2.5 μm 1/e2 diameter spot
and was aligned to the edge of the film to couple the 1490 nm (or 980 nm) pump laser
which was directly modulated with ~10 ms long pulses. The fluorescence was then
coupled back into the lensed fibre and then the emission component around 1550 nm
extracted with the pump power rejected by a chain of three 1490 nm (or 980 nm)/1550
nm wavelength division multiplexers (WDMs) providing in excess of 90 dB pump
rejection. Detection was accomplished with a connectorised 150 μm diameter InGaAs
diode and a fast, low noise Signal Recovery Inc. transimpedance amplifier. Data was
captured on a PC equipped with a 16 bit analog to digital converter, and custom
software enabled arbitrary amounts of trace averaging to remove noise [152]. The rise
57
and fall time of the complete light modulation and detection system was less than 20 μs
allowing accurate lifetime measurement even for sub-ms lifetimes. It should be noted
that the capability of recording the photoluminescence intensity during lifetime
measurement made the system usable in studying the relationship of PL versus pump
power.
Figure 3.7 Optical set-up for PL lifetime measurement of Er3+:As2S3 film [152].
Generally, the 1/e and intrinsic lifetimes are different, and the 1/e lifetime is always
pump intensity dependent. The intrinsic lifetime is approximated as the measured decay
rate of the PL once it decays >10-100x from its starting value. At this point, intensity
dependent effects become weak and the number of excited ions has significantly
reduced eliminating many of the ion-ion effects.
For the 1.8x1020 ions/cm3 Er3+ doped As2S3 films pumped at 1490 nm with 350
kW/cm2 intensity, the 1/e and intrinsic lifetime are 0.7 ms, and 0.9 ms, respectively. The
difference between these two lifetimes indicates the presence of ion-ion interactions.
The 1.55 µm PL decay curve of an Er3+:As2S3 film is shown in Figure 3.8(a). If the 1.55
µm emission is strong enough not to fall into the noise floor of an optical spectrum
analyser (OSA), replacing the photodiode with an OSA, the PL spectrum of the film can
be obtained. A 1.55 µm emission spectrum from a 600 nm thick erbium ion doped
As2S3 film is shown in Figure 3.8(b), where a broad emission band with peak centred at
1538 nm is clearly seen.
58
Figure 3.8 1.55 µm PL decay curve of an Er3+:As2S3 film (a), and a PL spectrum of 1.55
µm emission (b), from as-deposited erbium ion doped As2S3 films under excitation at 1490
nm.
The dependence of lifetime and PL intensity on pump intensity was investigated in
an Er3+:As2S3 film with 1.8x1020 ions/cm3 erbium ion concentration, and the results are
shown in Figure 3.9. From the results, the 1/e lifetime drops from 0.7 ms under an
excitation intensity of 10 kW/cm2 to around 0.6 ms when the excitation intensity
reaches 415 kW/cm2. This dependence of the 1/e lifetime on the pump power indicates
both ion-ion energy exchange effects and excited state processes are occurring when
high excitation power is involved. The occurrence of ion-ion energy exchange effects
may also be reflected in the trend of PL intensity increases with excitation power. If the
sample is not close to the saturation region, the PL output power would be expected to
be linear with pump power in the absence of energy exchange effects. The quadratic
trend relationship between PL intensity and pump power observed (Figure 3.9 (b))
instead of a linear relationship indicates that radiative and non-radiative ion-ion energy
exchange interactions are present in film. When measuring the PL emission with high
pump intensity, weak green light (~10-13 W from the OSA spectrum) was observed to be
emitted from the film edge with the naked eye in a dark environment, which is a clear
sign that up-conversion is occurring. To determine the exact processes participating
requires a much more thorough examination and will be reported later in Section 3.4.
59
Figure 3.9 1/e lifetime (a), and 1.55 µm emission intensity (b), versus excitation intensity at
1490 nm of an Er3+:As2S3 film after 2 hours 130 °C thermal treatment.
(b) Effects of “standard” thermal post-treatment practices on erbium ion doped
As2S3 films
Thermal annealing in rare-earth ion doped materials would be expected to reduce the
defects in the host material and also ‘activate’ the doped rare-earth ions provided the
solubility limit of the glass has not been exceeded. Several reports [129, 130, 217]
showed significant 1.5 µm PL intensity enhancement in rare-earth ion doped
chalcogenide glasses after thermal treatment (up to 40x in Fuchs’s report) [217]. Also a
longer lifetime of the 4I13/2 state of Er3+ was anticipated and has been observed [129,
130] after this treatment.
Lifetimes of the 4I13/2 state of Erbium ions in Er3+:As2S3 films (~2x1020 ions/cm3)
after thermal or light treatment were measured under two different excitation
wavelengths, 980 nm and 1490 nm, respectively. With the lifetime measurement set-up
described in Section 3.3.2(a), the error of lifetime measurement is about ±0.005 ms. The
results are shown in Table 3.2 with a pump intensity estimated at 200 kW/cm2. The
difference between 1/e and intrinsic radiative lifetime as noted earlier (see Section
3.3.2(a)) indicates the presence of some non-radiative decay, energy exchange
mechanism, or up-conversion process. It is clear that 1490 nm and 980 nm excitation
give similar results. This is somewhat surprising, given the 4I11/2 state would normally
has a lifetime similar to that of the 4I13/2 state in a low phonon energy host [46] and so
some “energy storage” effect would be expected that manifests as a longer 1550 nm
lifetime for 980 nm pumping [238]. The reasons for this are not clear, but one upside is
that this means that an accelerated decay is occurring from the 4I11/2 state may enhance
pumping efficiency for the higher inversion 980 nm pumping mechanism. There is little
difference between the results for thermal treatment and light treatment compared to the
60
as-deposited case, indicating that the long and medium range glass structure has little
impact on the rare-earth ion performance.
Table 3.2 Lifetimes of the 4I13/2 state of erbium ion in Er3+:As2S3 films after thermal or
light treatment.
Excitation @ 1470 nm Excitation @ 980 nm
Treatment of Er3+:As2S3 films 1/e lifetime
/ ms
Intrinsic
lifetime / ms
1/e lifetime
/ ms
Intrinsic
lifetime / ms
As-deposited 0.47 0.65 0.48 0.66
Thermal annealing
(130 °C, 24 hours under
vacuum)
0.5 0.69 0.56 0.72
Light annealing
(500-570 nm green light,
~3mW/cm2, 48 hours, in
ambient atmosphere)
0.49 0.64 0.49 0.65
(c) Propagation loss and absorption band of erbium ion doped As2S3 film
Measurements of film loss and Er3+ absorption band (range from 1450 nm to 1650 nm)
were performed using a Metricon 2010 prism coupler, tunable laser sources, and an
InGaAs camera. The scattering streaks from the TE fundamental slab guided mode was
imaged and captured using a cooled high sensitivity InGaAs camera at wavelengths
between 1450 and 1650 nm and then the propagation loss was calculated using custom
image processing software by analysing the decay of the normalized and background
corrected scattering streak versus distance. A representative beam streak is shown in
Figure 3.10(a). The resulting optical loss of the 1.8x1020 ions/cm3 Er3+ doped As2S3 film
(after 2 hours 130 °C thermal treatment) is shown in Figure 3.10(b).
61
Figure 3.10 Beam streak in an Er3+:As2S3 film at 1500 nm, light entered on the right side
and exited on the left at the edge of the wafer (a), and the absorption curves (b) of an Er3+
doped As2S3 film.
The absorption band peak due to the 4I15/2→4I13/2 transition of Er3+ lay at 1538 nm,
which is slightly red-shifted from the 1532 nm usually observed in oxide glass hosts
[239-241], and blue-shifted from that observed in Er2O3 crystal [155]. This red-shift in
absorption band is normal in chalcogenide glass host materials [46, 102, 242] and could
be explained by the nephel-auxetic effect (‘electron cloud expanding’), that is, rare-
earth ion radiative absorption and emission tend to red-shift slightly with increasing
covalency of the lattice [39]. The obtained red-shift of the absorption curve and weaker
absorption tail at shorter wavelength indicated a longer excitation wavelength for
optimum pumping is required than for the normal 1480 nm oxide host pump lasers.
3.4 Up-conversion properties of erbium ion doped As2S3 film
Detailed investigation of the up-conversion processes present in the material enable a
better prediction of the emission properties and potentially available optical gain, so
attempts were made to quantify the various up-conversion products, and their
dependencies on pump intensity and doping concentration. The starting point was
measurement of the emission spectrum across a wide range (400-1700 nm) that is able
to cover all necessary emissions and up-conversion bands. Due to the wide range, two
different detectors (Ocean Optics HR4000 spectrometer and an Ando AQ6317 OSA)
were employed in the optical set-up. For a better accuracy, these two detectors were
supposed to work at the same time, corresponding to the same excitation intensity. The
scheme of the set-up for this measurement is shown in Figure 3.11. In the measurement,
62
erbium ion doped As2S3 film is pumped with 1490 nm excitation through a clean,
cleaved edge via a lensed silica fibre. Emission captured by the same lensed fibre is
recorded using the OSA for the 700-1700 nm range. Emission is also captured by a 250
micron core silica fibre from the top of the excited film edge, then passes through a
short-pass filter placed in a pair of collimators and finally is analysed using an Ocean
Optics HR4000 spectrometer to cover the 400-1000 nm range. As the OSA and Ocean
Optics spectrometer have a spectral overlap across 600-1000 nm, and given that the
OSA has power calibration and normalization, then an Ocean Optics white light source
is used as a transfer reference to cross-calibrate power and level the response spectrum
of the Ocean Optics spectrometer.
Figure 3.11 Optical set-up for wideband erbium ion emission in an Er3+:As2S3 film from
visible to NIR.
The combined spectra collected by both detectors versus pump power are shown in
Figure 3.12. The PL band at 1540 nm arises from the 4I13/2→4I15/2 emission. The sharp
peak at 1490 nm is the remnant pump laser, while the other bands around 980 nm, 805
nm, 670 nm, 540 nm and 520 nm correspond to the emission from the 4I11/2, 4I9/2,
4F9/2,
4S3/2 and the 2H11/2 energy states to the ground state, respectively, and potentially some
higher inter-level transitions. Possible radiative transitions of Er3+ are shown in Table
3.3, including the related energy levels, wavelength and calculated spontaneous
emission probabilities based on 98(0.15Ga2S3-0.85GeS2)-2Er2S3 (mol%) glass in ref
[146].
63
Figure 3.12 Emission spectrum of an Er3+-doped As2S3 film (after 2 hours 130 °C thermal
treatment) pumping at 1490 nm from visible to NIR versus pump power.
Table 3.3 Possible radiative transitions of Er3+ and calculated spontaneous emission
probabilities [A(J;J׳) (s-1)] based on 98(0.15Ga2S3-0.85GeS2)-2Er2S3 glass in ref [146].
Emission wavelength /
µm
Transition A(J;J׳) (s-1)
Initial state Final state
0.525 2H11/2 4I15/2 14370
0.546 4S3/2 4I15/2 2371
0.656 4F9/2 4I15/2 4718
0.78 2H11/2 4I13/2 --
0.803 4I9/2 4I15/2 513
0.848 4S3/2 4I13/2 964
0.98 4I11/2 4I15/2 1029
1.05 2H11/2 4I11/2 --
1.15 4F9/2 4I13/2 444
1.25 4S3/2 4I11/2 77
1.45 2H11/2 4I9/2 --
1.53 4I13/2 4I15/2 508
1.7 4I9/2 4I13/2 94
1.7 4S3/2 4I9/2 156
2.0 4F9/2 4I11/2 397
2.7 4I11/2 4I13/2 99
3.6 4F9/2 4I9/2 73
4.5 4I9/2 4I11/2 6
-- Data unavailable
64
The erbium ion concentration in this sample is relatively high, at 1.8x1020 ions/cm3,
leading to a drop in the average distance between ions, and therefore to enhanced ion-
to-ion interactions due to the r6 (r is the distance between neighbour ions) dependence
[149]. With high excitation intensity, ions in the 4I13/2 metastable state with a relatively
long lifetime of 0.9 ms have a high probability of promotion to the 4I9/2 energy state via
ETU, ESA and cooperative up-conversion processes (see Section 2.3.2). At this point
erbium ions may experience several different transitions. Firstly, erbium ion can
radiatively decay to the ground state with corresponding emission of 805 nm photons.
Ions in the 4I9/2 state may also decay to the 4I11/2 state by multiphonon relaxation or
radiative decay with a longer wavelength emission that could not be recorded with the
OSA. Ions in the 4I11/2 energy state explained the 980 nm emission in the spectra. An
intrinsic lifetime of 0.5 ms was measured for the 4I9/2 excited state, and this lifetime
allows the possibility that ions in the 4I9/2 excited state receive energy from another ion
in the metastable state 4I13/2 via an ETU process and elevate to the higher 2H11/2 energy
level, which results in the 520 nm emission. The energy difference between the 2H11/2
and the 4S3/2 levels is small, therefore ions at the 2H11/2 level can non-radiatively decay
to the 4S3/2 level quickly with the help of host phonons. The small intensity 540 nm
emission is assigned to the radiative decay from the 4S3/2 state to the 4I15/2 ground state.
The emission band at 670nm could be explained by the radiative decay from the 4F9/2
state to the ground state, on which excited states accumulate by phonon assisted non-
radiative decay from the 4S3/2 upper energy level. It is worth to note that emissions of
540 nm and 520 nm are located in the bandgap of the As2S3 host material, which implies
emissions in this wavelength range would suffer extremely high absorption by the host.
As a result of this, the 540 nm and 520 nm emission intensity in these spectra may be
much lower in intensity than the real value. Also in chalcogenide glass hosts, the
intensity of the 980 nm emission is much smaller than the 805 nm emission, which is
different from the results in most oxide glasses [175]. This could be explained as
follows, based on the erbium ion energy level system (see Fig 2.1), the 980 nm emission
arises from the radiative decay from the 4I11/2 state to the ground state, and the
population in the 4I11/2 state is due to the decay of ions from the 4I9/2 state to the 4I11/2
state. Due to the low phonon energy of chalcogenide glass hosts, the probability of this
decay is rather lower than that in high phonon energy hosts like oxide glasses, which in
turn lead to a weak emission at 980 nm.
It is worth noting that whilst the high intensity emissions at 980 and 805 nm are
present, their power spectral density is still at least an order of magnitude lower than
65
that at 1540 nm. These energy transfer processes are then causing some loss of pump
efficiency and to an extent depopulating the metastable level: 4I13/2, the first excited state.
As mentioned previously (Section 2.3), all the up-conversion emissions observed are
caused by a combination of the up-conversion processes, such as energy transfer up-
conversion (APTE or ETU), excited state absorption (ESA), etc. Besides the first-order
up-conversion from the metastable state, second-order cooperative up-conversion
processes are also occurred, which lead to the observed green emission at 520 and 540
nm, due to the 2H11/2→4I15/2 and 4S3/2→
4I15/2 transitions, respectively.
Although all these three processes lead to a similar short wavelength emission, there
are significant differences between them. For example, in ETU, two ions in the first
excited state (4I13/2) interact to populate the 4I9/2 state, and the 4I11/2 state is populated due
to rapid nonradiative or radiative decay from the 4I9/2 state. Therefore, the rate of up-
conversion depends quadratically on the concentration of ions in the 4I13/2 state (i.e.
dopant concentration both for the population and the degree of interaction via
separation), and therefore also quadratically on the pump power. In contrast, excited
state absorption (ESA) involved only one excited ion, this process is independent of
erbium ion concentration but only related with pump intensity.
In order to determine which of these mechanisms is responsible for the observed up-
conversion luminescence, one popular method is to measure the luminescence intensity
of the different bands as a function of pump power. With the data analysed
quantitatively, the up-conversion coefficient and ESA cross-section could be determined
[175].
To accomplish this task, the optical set-up shown in Figure 3.11 was modified. An
optical attenuator was inserted in the pump laser line to control precisely the excitation
power and enable the power to be tuned over a big range (>30 dB) as was required to
obtain a good fit [175]. In the lensed fibre output side, a power meter with -110 dBm
sensitivity was used instead of the OSA to record the power of 1.5 µm emission, and
two more 1490/1550 nm WDMs were applied to remove completely the residual pump
power. Narrow band-pass optical filters and a Si detector with a low noise trans-
impedance amplifier having gain of 109 V/A (volts/amps), and a lock-in amplifier
system were employed to record the pump power dependence of emission from the up-
conversion processes. The experimental results are plotted in Figure 3.13.
66
Figure 3.13 Luminescence intensity at 1.5 µm (a), 800 nm (b)and 980 nm (c) vs. excitation
power and 1.5 µm emission of an erbium ion doped As2S3 film, on log10 scales.
The 1.5 µm emission intensity versus pump power at 1490 nm of a 1.8x1020 /cm3
erbium ion doped As2S3 film is plotted in Figure 3.13 (a), in log10 scale. A linear trend
is found between 1.5 µm emission and pump power, but the slope is 0.887 from the
fitted line, implying processes such as energy transfer and up-conversion, that can
reduce 1.5 µm photon emission yield, are occurring. PL intensity at 800 nm and 980 nm
versus that at 1.5 µm are shown in Figure 3.13 (b) and (c), respectively, and both curves
can be linear fitted with slopes of 1.87, and 1.56, respectively. Both slopes are smaller
than 2 as might be expected if no non-radiative decay and/or second-order up-
conversion occurs. Due to the long lifetime of the 4I11/2 state (intrinsic lifetime of 0.84
ms) in this sample, ions in the 4I11/2 state may be promoted to the 4F9/2 state by photon
absorption or energy transfer, which depletes the population of the 4I11/2 state, resulting
in a slope less than 2. Similar phenomena can occur at 800 nm. With a long 4I9/2 state
lifetime of 0.59 ms, ions could be further excited to the 4S3/2 state from which the 540
nm emission arise to the ground state.
Since the ions can reach the 4S3/2 state via second-order up-conversion or a 3-photon
process, then it is also possible to decay from the 4S3/2 state to the 4I13/2 state and give
out emission around 800 nm. The zoomed in 800 nm emission spectrum shown in
Figure 3.14 supports this assertion. The broad 800 nm emission band can be divided
into two bands almost overlapping each other. The one centred at 810 nm is proposed to
be associated with the 4I9/2 state to the ground state transition, whilst the other one
located at 830 nm is thought to arise from the decay between the 4S3/2 state and the 4I13/2
state.
67
Figure 3.14 Emission spectrum at 800 nm of an erbium ion doped As2S3 film of 1.8x1020
ions/cm3 Er3+ concentration after 2 hours 130 °C thermal treatment.
In this research, even with the maximum pump power the diode laser can reach,
ASE saturation was not evident, and therefore it was difficult to calculate the up-
conversion efficiency by modelling with rate equations under steady-state conditions
[175]. A result showing such saturation behaviour was achieved in an erbium ion doped
TeO2 reference film with ~2.2x1020 ions/cm3 concentration, but the reasons behind the
different behaviour in erbium ion doped As2S3 films are still unclear.
3.5 High temperature thermal post-treatment on erbium ion
doped As2S3 film
It is well known that high temperature thermal treatment often yields positive effects on
rare-earth ion doped materials. With the energy from thermal treatment, dopants are
able to ‘re-bond’ and be ‘active’ and better incorporate themselves into the host
material. But at the same time, given the known low solubility of erbium ion in bulk
As2S3 [118, 197], the question naturally arises as to whether the thermal or light post-
treatment process used to restore the film towards the bulk state causes erbium ion
precipitation.
Therefore, the question of whether thermal treatment up to and above Tg helps or
hinders is relevant, and gives insight into the effects of the “standard” 130 °C 24 hours
thermal treatment typically applied to As2S3 films. Based on previous experiments in
the ANU’s Laser Physics Centre, thermal treatment of chalcogenide films at
68
temperatures beyond the Tg, will not damage the film directly in terms of film quality
observed in dark field microscopy, provided the thermal treatment occurs in a vacuum
below Tg and an inert atmosphere above Tg.
Several small pieces (10x10 mm) of erbium ion doped As2S3 films of 0.6 at% Er3+
concentration (1.8x1020 ions/cm3) were cleaved and prepared for different thermal
treatments. One of the Er3+:As2S3 films was also measured as deposited and then
subjected to the standard 130 °C 24 hour thermal treatment under vacuum and
remeasured for comparison. Figure 3.15 presents the results.
Figure 3.15 Results of thermal treatment on 0.6 at% Er3+ doped As2S3 films. Samples were
thermally treated for 1 hour at temperature indicated, except 130 °C data point which is
for 24 hours under vacuum.
There are two interesting features from Figure 3.15. Firstly, it is apparent that as the
thermal post-treatment temperature rises, the lifetime of the active erbium ions is
increasing by up to ~50%. This indicates that the erbium ions are not optimally
incorporated into the glass matrix after deposition nor after the standard thermal
treatment. Secondly, the PL intensity drops markedly as the thermal post-treatment
temperature exceeded ~130 °C and drops rapidly to zero as the glass Tg (~180 °C) is
exceeded. For reference purposes, the sample measured before and after the 130 °C 24
hour thermal post-treatment under vacuum exhibited about a 30% reduction in PL
intensity at the 1.8x1020 ions/cm3 doping used. It is not clear how much of the drop in
PL intensity due to the 2 hours thermal post-treatment at 130 °C is attributable to
erbium ion diffusion and clustering as opposed to changes induced by the large
structural rearrangement of the host glass, but the reduction is tolerable given the known
benefits to the glass stability and properties [213].
69
3.6 Concentration effects in erbium ion doped As2S3 films
Er3+:As2S3 films with different erbium ion concentrations were deposited using co-
thermal evaporation. Bulk As2Se3 (5N) and metallic erbium were used as starting
materials. The deposition details were described previously (see Section 3.2.3), and the
only difference was the evaporation temperature of erbium metal source. During
deposition, the erbium metal target temperature was decreased from 1070 °C to 1030 °C
and 990 °C, respectively, to drop the erbium metal evaporation rate to a half and a
quarter of its previous value of 0.001 nm/s (see Section 3.2.3), whilst the power and rate
for As2S3 target was kept as previously at 0.1 nm/s. By doing so, samples with 0.15 at%,
0.3 at% and 0.6 at% erbium ion doped concentrations were fabricated. It is worth
pointing out that with the current deposition system, the erbium ion concentration in the
as-deposited As2S3 films could not go lower than 0.15 at% because the erbium metal
target evaporation rate here was at the measurement limit.
All the as-deposited wafers were coated with a 120 nm thick SU-8 layer to protect
from contamination, and were subjected to 130 °C thermal post-treatment in a vacuum
for 24 hours, and then two days of light post-treatment with 500-570 nm green light
(~3mW/cm2) in ambient atmosphere to try to ensure the Er3+:As2S3 films were fully
relaxed toward the condition of their bulk counterpart.
Lifetime and PL intensity versus excitation intensity for different concentrations of
Er3+ in these Er3+:As2S3 films were measured and are plotted in Figure 3.16. From the
results, the 0.15 at% sample has the longest lifetime of 1.2 ms, the lifetime of the 0.3
at% sample drops to 0.9 ms, while the highest doped sample has the shortest lifetime of
about 0.5 ms. Also lifetimes are relatively independent of pump power: taking the 0.15
at% sample as an example, the 1/e lifetime only fell from 1.23 ms (excitation intensity
of 50 kW/cm2) to 1.17 ms (excitation intensity of 1100 kW/cm2), and the PL decay
curve showed a single exponential nature. Conversely, with increasing erbium ion
concentration, the lifetime drops significantly, indicating strong ion-ion interaction
induced photoluminescence quenching occurs. In the PL intensity curves in Figure
3.16(b), the PL intensity of all three samples increases with excitation intensity, but in a
quadratic way rather than linear way, especially at high excitation intensity, indicating
that even in the 0.15 at% doped sample ion-ion interaction occurs as well. Also, the PL
intensity of the 0.3 at% sample is less than double that of the value of the 0.15 at%
sample, implying the PL yield per erbium ion drops at higher concentrations. In the
most heavily doped 0.6 at% sample (1.8x1020 ions/cm3), the film edge is easily damaged
70
with high excitation power. This phenomenon is assumed to be due to the higher
absorption with the greater doping, coupled with a higher degree of ion-ion interaction
and energy migration, meaning a large amount of the pump energy absorbed is
transferred to defects or phonons, which in turn damages the film edge by local
temperature elevation. Therefore, there are no data recorded at excitation intensities
over 400 kW/cm2 for the 0.6 at% film.
Figure 3.16 1/e lifetimes (a), and PL intensities (b), of Er3+:As2S3 films with different
erbium ion concentrations.
Again, high temperature thermal post-treatment was applied to the 0.15 at% sample
for 1 hour at each fixed temperature. The high temperature thermal post-treatment
results are shown in Figure 3.17. With higher thermal post-treatment temperatures, the
lifetime increases from 1.23 ms up to 1.88 ms after thermal post-treatment at 280 °C,
while the PL intensity drops to low values indicating perhaps erbium ion diffusion and
precipitation. These results provide conclusive evidence that the optimum additive of
erbium ion in As2S3 is < 0.15 at%.
Figure 3.17 Results of thermal post-treatment of 0.15 at% Er3+ doped As2S3 films. Samples
are post-treated for 1 hour at the temperature indicated, except for the 130 °C data point
which is for 24 hours under vacuum.
71
3.7 Conclusion
This chapter reported the fabrication of erbium ion doped As2S3 films by co-thermal
evaporation. Both structural and optical properties of the obtained films were
investigated in detail. Wideband emission spectra of an Er3+:As2S3 film were measured.
Emissions due to up-conversion and second-order up-conversion were characterised.
Influence of >Tg thermal post-treatment on lifetime and PL intensity in erbium ion
doped As2S3 film was studied and analysed. By comparing the performance of samples
with different erbium ion concentrations, it is found that the optimal erbium ion
concentration observed to date in an As2S3 film made using thermal evaporation is
about 0.15 at%. However, it was also pointed out the solubility limit of erbium ion in
As2S3 is below this value.
72
73
Chapter 4
Erbium ion doped As2S3 waveguide amplifiers
This chapter reports the fabrication of waveguide amplifiers based on erbium ion doped
As2S3 films using photolithography and reactive ion etching, as well as the performance
of the obtained waveguide amplifiers. Optical waveguide theory is briefly discussed
first, and the waveguide design and simulation for the devices are covered. Waveguide
fabrication is briefly described and then the propagation losses of the obtained erbium
ion doped waveguides as well as the absorption properties of erbium ions in the
waveguides, are characterised. Excitation experiments are finally performed to study the
amplification property of the waveguides.
4.1 Basic theory and simulation methods
4.1.1 Maxwell’s equations for optical waveguides
It is well known that the propagation of light can be described by Maxwell’s equations
[243] as follows:
∇ ∙ 𝐄 =𝜌
𝜀0 (4.1)
∇ ∙ 𝐁 = 0 (4.2)
∇ × 𝐄 = −𝜕𝐁
𝜕𝑡 (4.3)
∇ × 𝐁 = 𝜇0 (𝐉 + 𝜀0𝜕𝐄
𝜕𝑡) (4.4)
where E and B are the electric and magnetic fields respectively, ε0 is the free space
permittivity, μ0 is the free space permeability. ρ represents electric charge density and J
is the electric current density. If the light is propagating in a pure dielectric medium, (i.e.
a medium with no conductivity or charges), no magnetic effects are present, and in the
further case of an isotropic and linear system, then Maxwell’s equations reduce to:
∇ × 𝐄 = −𝜇0𝜕𝐇
𝜕𝑡 (4.5)
∇ × 𝐇 = 𝜀0𝑛2 𝜕𝐄
𝜕𝑡 (4.6)
74
where 𝐇 =1
𝜇0𝐁, and the n is the refractive index distribution of the propagation media.
If the medium is optically inhomogeneous its properties are position-dependent, n=n(r).
The wave equations for E and H can be derived as [244]:
∇2𝐄 + ∇(1
𝑛2∇𝑛2𝐄) − 𝜀0𝜇0𝑛
2 𝜕2𝐄
𝜕2𝑡= 0 (4.7)
𝛻2𝑯+1
𝑛2𝛻𝑛2 × (𝛻 × 𝑯) − 𝜀0𝜇0𝑛
2 𝜕2𝑯
𝜕2𝑡= 0 (4.8)
Normally, these two full vectorial equations describe all coupled components of the
electric and magnetic fields, and each component cannot be reduced to a scalar
equation. Although the dielectric medium for amplification may encounter intensity-
dependent refractive index, for the characterization of light propagation in the
waveguide, the numerical solutions consider linear refractive index. In many structures,
for example the waveguides used in this thesis, the lack of cylindrical symmetry
prevents analytical solutions, thus numerical methods must be used to calculate the
eigenvalues of the eigenmode equations which describe the propagating modes of the
waveguide.
4.1.2 Simulation of optical waveguide
Accurate modelling of waveguide-based devices plays several significant roles in the
advancement of optical devices, including optimization of current designs, shortening of
the design and fabrication cycle, and evaluation of new device concepts. So far,
modelling techniques for optical guided-wave propagation can be divided into two
groups: time-harmonic (e.g. monochromatic continuous wave (CW) operation) and
general time-dependent (e.g. pulsed operation) [245]. Excepting short pulse high
intensity cases, where significant nonlinear effects occur, time-harmonic methods are
usually accurate tools when implemented to an appropriate level of exactness.
In this work, to complete the design of waveguide amplifiers, modal analysis was
essential. Factors such as the number of modes supported by the waveguide, their
propagation constants, mode intensity distribution, modal excitation in response to a
given input field, and the overlap between the mode distribution and the erbium ion
doped area especially, can readily be extracted using modal analysis. Different
simulation methods, for example the finite difference method, finite element method,
mode-matching technique, beam propagation method (BPM), method of lines (MoL),
75
coupled mode theory, and the finite-difference time-domain method (FDTD), are
commonly used to compute the electromagnetic modes of waveguides [245].
The key objective for mode solvers is to determine the value of the propagation
constant and the corresponding mode field distribution for each mode the structure
supports in a given waveguide cross-section at a specified operating wavelength. The
finite difference and finite element methods (based on variational principles) have found
wide acceptance in the community with demonstrated accuracy even for high refractive
index contrast full vectorial cases. They possess advantages such as arbitrary structure,
arbitrary shape of waveguides, graded meshes, and numerically stable solution methods
for this method are readily available [246]. Detailed discussion on these two methods
can be found in [245-248].
4.2 Waveguide design
4.2.1 Waveguide geometry
Planar waveguide structures are primarily and widely used to integrate photonic
technologies to the chip scale and gain the same sorts of benefits enjoyed by the silicon
integrated circuit industry. Compared to fibres, they enable multi-functional multi-
material integrated devices, and also combine higher light confinement and high
refractive index contrast. The most common types of waveguide are the buried channel
and ridge based devices, which are shown in Figure 4.1(a) and (b), respectively. A
typical structure of a buried channel device uses three different refractive index
materials and can be described with two geometric parameters, Width and Height. A
generalised ridge waveguide is composed of four different materials and can be
described by three geometric parameters, Width, Height and Etch depth as shown in
Figure 4.1(b). Normally, buried channel devices have better optical confinement and
thus have smaller bending radii and therefore advantages in realizing high density
photonic circuits. Ridge waveguides have lower propagation losses due to reduced
interaction of the mode with the waveguide sidewall and offer fabrication opportunities
unavailable to buried channel devices [249].
76
Figure 4.1 Structures of: buried channel waveguide (a), and ridge waveguide (b). In this
figure, W is the width, H is the height and D represents the etch depth of waveguide.
In this work, there are particular considerations that make the ridge waveguide
advantageous. These stem mainly from the desire to make the waveguides by
lithography and dry etching where one material cannot easily be etched as this is an
established route to low loss devices in chalcogenide glasses [13, 250].
(a) Waveguide width and etch depth
A structure that is as simple and easily achievable as possible is preferred. In waveguide
design, the total film thickness should not be too big, which will bring issues in film
deposition, thermal treatment and film etching, therefore based on experience with
annealing of chalcogenide glass films, the film thickness was fixed at around 1.3 µm.
Modal effective refractive index was studied as a function of etch depth for
waveguides of different widths (1, 2 and 3 µm), and the results are shown in Figure 4.2.
In simulation, the refractive index of the thermally oxidized silicon (TOX) substrate is
1.445, the core (take As2S3 as an example here and fix n2=n3) is 2.41 with an air
cladding (n1=1). In a 1 µm width waveguide, the etch depth needs to be above 300 nm
to stop the TE0 mode leaking into the slab, and 550 nm for the TM0 mode. In the wider
waveguides, 2 and 3 µm, smaller etch depths are required to confine properly the modes
but high order modes start to be guided. In the 2 µm waveguide, if the etch depth is
controlled in a range of 350 nm to 550 nm, single mode propagation at 1550 nm is
feasible, which is preferable for our application. For the 3 µm waveguide, high order
modes are guided at almost every etch depth. Therefore, in the following part, attention
is directed primarily to the 2 µm wide waveguide.
77
Figure 4.2 Etch depth versus effective refractive index of supporting modes for: 1 µm (a);
2 µm (b) and 3 µm (c) As2S3 waveguide.
(b) Bend and junction loss
The bending loss has to be taken into consideration during design of integrated devices.
A large bend radius would make the device less compact, while too small a bend radius
can introduce considerable extra loss. Therefore, a bending loss simulation was
completed to quantify the minimum bend radius the structure could reach without
causing significant loss. Commercial software Rsoft-FemSim was employed here to
undertake this task. In simulations, a 0.5 µm perfect matching layer (PML) was applied
at the domain edge and leaky modes were enabled in the calculation which relied on a
conformal transform method to simulate the bend [251]. The imaginary part of the
effective refractive index was used to calculate the bending loss using the following
formula:
𝑙𝑜𝑠𝑠 = 4𝜋 × 𝑛𝑖𝑚 × 4.3429 𝜆⁄ (4.9)
where nim is the imaginary part of the effective refractive index, and λ is the wavelength.
In Figure 4.3, both the fundamental TE0 and TM0 mode bending loss in the 2 µm wide
structure are shown.
Figure 4.3 The bending loss of the fundamental mode: TE0 mode (a) and TM0 (b) mode in
the 2 µm wide structure.
78
From the results in Figure 4.3, bending losses of both TE and TM modes increase
with decreasing bend radius in a roughly exponential way. For any radius above 400
µm, the bending loss is less than 0.01 dB per turn for both modes, which is negligible in
a real device. Similar bending loss curves were calculated for the 3 µm waveguide, and
the results are shown in Figure 4.4. Therefore, allowing a safety factor, a 400 µm radius
could be set as a minimum threshold to the occurrence of significant bending loss. It
should also be noted that, the bending loss here is purely through radiation. There is also
the mode shape shift introduced by the bending structure which causes a mode overlap
loss at the entrance and exit of any bends which is referred as junction loss. The
junction loss for 2 µm and 3 µm waveguides were calculated and are shown in Figure
4.5.
Figure 4.4 The bending loss of the fundamental mode: TE0 (a) and TM0 mode (b) in the 3
µm width structure.
Figure 4.5 The junction loss in: 2 µm (a) and 3 µm (b) width structure.
Junction loss drops almost exponentially with the increase of bend radius. For the 2
µm waveguide, junction loss is below 0.01 dB when the bend radius is over 400 µm for
the TM mode, and over 600 µm for the TE mode. In a 3 µm waveguide, due to the
79
bigger mode area the junction loss is little bigger than 2 µm waveguide with the same
bend radius, but still down to ~0.03 dB when the bend radius exceeds 600 µm for both
TE and TM mode.
(c) Hybrid integrated waveguide structure
A significant barrier exists towards fabricating the designed waveguides in rare-earth
ion doped materials, namely that there are no volatile halides or hydrides that can be
used to transport the etched rare-earth atoms [252]. The involatile erbium etch products
cause roughness as they tend to form micro-masks which then result in rough etched
surfaces and sidewalls. For example, Figure 4.6 shows the results of etching pure TeO2
and Er3+:TeO2 using an optimised hydrogen/methane/argon mix in an inductively
coupled plasma (ICP) reactive ion etching (RIE) system [152]. The degradation of the
etching quality is unacceptable, and an alternative solution is required. To avoid this
issue, a hybrid integrated waveguide structure is proposed. In a hybrid integrated
waveguide structure, more than one materials are combined to achieve more flexibility
and better performance.
Figure 4.6 Images of etched: un-doped TeO2 (a) and Er3+:TeO2 (b) waveguides [152].
Here, a strip loaded waveguide was investigated comprising an erbium ion doped
As2S3 slab and an un-doped As2S3 loading strip. In this structure, H=1350 nm total
thickness was used, with h =500 nm etch depth and 50 nm pure As2S3 left as a buffer,
and the Er3+ doped As2S3 layer is 800 nm thick. A small overlap reduction between the
mode and the doped area is the price to avoid the issue of etching erbium ion doped
As2S3 directly. The cross-section of the finalised rib waveguide structure is shown in
Figure 4.7, and the TOX refers to thermally oxidized silicon substrate.
80
Figure 4.7 The cross-section of the finalised ridge waveguide structure to avoid etching
Er3+:As2S3 directly, and TOX is the thermally oxidized silicon substrate.
With this finalised ridge waveguide structure, the properties of the propagation
modes in this waveguide were simulated in the following section.
4.2.2 Mode properties of the hybrid integrated waveguide structure
In the waveguide design, besides parameters such as effective refractive index, mode
area and the mode distribution, two more specific parameters are also of interest. Firstly,
the overlap between the optical mode distribution and the erbium ion doped area (see
Figure 4.7) is important. This overlap is less than unity because of field penetration into
the un-doped “strip loading” layer in the two layer structure, which in turn decreases
gain and the pump efficiency. Therefore, a structure with a large overlap is preferable.
Secondly, the coupling loss from an input fibre into the waveguide is of considerable
interest and may be calculated from the mode overlap formulism [244].
Simulations were performed using a public domain Matlab finite difference code
written by Fallahkhair et al., 2008 [253]. The Matlab finite difference mode solver is a
full-vector simulation based on the transverse magnetic field components. It is versatile
and reprogrammable to allow custom properties to be calculated. Waveguides of the
structure shown in Figure 4.7 were modelled with variable widths of 1 µm, 2 µm, 3 µm
and 4 µm, respectively, matching the widths on the standard waveguide test mask used
at the Laser Physics Centre, in the Australian National University. In calculations, a grid
size of 20 nm was applied with a 10 µm × 5.35 µm window, and the wavelength used in
calculations was 1.55 µm. In the simulation, a perfectly matched layer (PML) was
applied at the domain edge. The resulting TE and TM mode distributions for different
waveguide widths are shown in Figure 4.8. Effective refractive index, mode area
(defined as in [254]), overlap between optical mode distribution and the erbium ion
doped area and coupling loss from lensed fibre (2.5 µm of 1/e2 diameter with 14 μm
81
working distance) to waveguide were obtained from calculation and are summarised in
Table 4.1.
Figure 4.8 Simulations of the TE and TM mode distributions for waveguide structure
depicted in Figure 4.7 with different waveguide widths: 1 µm (a); 2 µm (b); 3 µm (c) 4 µm
(d). The horizontal dotted black line in each figure represents the structure of waveguide.
Note the colour bar scale represents major electric field component normalised to unity
peak value.
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Table 4.1 Simulation results of mode effective refractive index, mode area, the overlap
between mode and Er3+ doped area and fibre-waveguide coupling loss for the waveguide
structure depicted in Figure 4.7.
TE0 TM0
Waveguide
width / µm
Mode
effective
refractive
index
Mode
area /
µm2
Overlap
between
mode
and Er3+
doped
area / %
Coupling
loss / dB
Mode
effective
refractive
index
Mode
area /
µm2
Overlap
between
mode
and Er3+
doped
area / %
Coupling
loss / dB
4 2.355 3.31 65.5 1.76 2.342 2.97 67.9 1.90
3 2.351 2.64 66.5 1.56 2.338 2.34 68.4 1.68
2 2.342 2.04 70.0 1.61 2.329 1.75 70.1 1.78
1 2.319 1.97 85.0 1.92 2.302 1.75 81.3 2.09
Based on the results shown in Table 4.1 and Fig 4.8, wider waveguides have bigger
mode area (unsurprisingly but in a sub-linear manner) and less overlap of the mode with
the erbium ion doped area. They also, however, have lower interaction with the
sidewalls which should produce lower scattering losses. The 1 µm and 2 µm
waveguides both have significantly better overlap than the other widths, but the 1 µm
waveguide has the highest coupling loss due to the mismatch between waveguide mode
area and the output spot size of the lensed fibre. Narrow waveguides also have increased
sidewall overlap and will suffer more bending losses as discussed in the previous
section (see Section 4.2.1), and so the best choice is the 2 µm wide waveguide.
4.3 Simulation of amplification performance of erbium ion
doped As2S3 waveguide amplifier
4.3.1 Rate equation for an erbium ion doped waveguide system
The dynamics of energy level populations in laser gain media are often modelled using
a system of rate equations. These are differential equations, describing the temporal
evolution of level populations under the influence of optically induced and non-
radiative transitions. To predict the performance of an Er3+:As2S3 waveguide amplifier,
simulation based on rate equations was performed. Laser diodes working at 980 and
1480 nm are both practically used for Er3+ doped fibre amplifiers, and it was reported
83
that 980 nm pumping was shown to have better gain characteristics and lower noise
figure [255, 256]. In chalcogenide glasses, due to the low phonon energy, the lifetime of
the 4I11/2 state is long (intrinsic lifetime of 0.84 ms, see Section 3.4), resulting in
significant up-conversion processes when pumping at 980 nm. To avoid this, in this
thesis, the amplification experiment was carried out with 1480 nm excitation. With 1480
nm excitation, the erbium ion works as a quasi-three level energy system (see Section
2.1). In this case, the stark levels of the 4I13/2 state are divided into two energy groups,
the upper energy level and lower energy level. Erbium ions in the ground state absorb
pump photons and are promoted into the upper part of the 4I13/2 excited state. These then
rapidly decay non-radiatively due to interactions with the thermal phonon bath and
through electron-electron energy loss processes to the metastable state, the lower energy
part of the 4I13/2 level. With the ion accumulation in the lower 4I13/2 energy level, a
population inversion between it and the ground state 4I15/2 is formed. During this
process, basic up-conversion effects are taken into consideration in the model, therefore,
the higher excited state 4I9/2 is also in the picture. As the pump energy is approximately
equal to the 4I9/2 - 4I13/2 energy difference, ions in the lower 4I13/2 energy level may get
further excited (by absorbing another pump photon (ESA) or accepting energy from
neighbouring excited ions (co-operative upconversion)) and be promoted to the 4I9/2
energy state. Further processes can also occur from this state but were not included in
the model. A detailed energy level diagram of erbium ion in a glassy host and the main
processes modelled with 1480 nm excitation are shown in Figure 4.9.
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Figure 4.9 Energy level diagram of erbium ion in a glassy host with all potential energy
transfer mechanisms. In this diagram, Rij and Wij are the pump and signal stimulated rates,
the Aij are the nonradiative rates from level i to j, A21 is the fluorescence rate, ESA is
excited state absorption, Cup is the homogeneous upconversion coefficient, and the C14 is
cross-relaxation coefficient. R13 and W21 represent the primary pump and primary PL,
respectively.
To describe ion state populations in an erbium ion doped glass, a quasi-three level
energy system is used as described above, the rate equations would normally be written
as follows:
𝜕𝑁1
𝜕𝑡= −𝑊12𝑁1 − 𝑅13𝑁1 + 𝑅31𝑁3 + 𝐴21𝑁2 +𝑊21𝑁2 + 𝐶𝑢𝑝𝑁2
2 − 𝐶14𝑁1𝑁4 (4.10)
𝜕𝑁2
𝜕𝑡= 𝑊12𝑁1 − 𝐴21𝑁2 −𝑊21𝑁2 + 𝐴32𝑁3 − 2𝐶𝑢𝑝𝑁2
2 + 2𝐶14𝑁1𝑁4 − 𝑅24𝐸𝑆𝐴𝑁2 (4.11)
𝜕𝑁3
𝜕𝑡= 𝑅13𝑁1 − 𝑅31𝑁3 − 𝐴32𝑁3 + 𝐴43𝑁4 (4.12)
𝑁1 + 𝑁2 + 𝑁3 +𝑁4 = 𝑁𝐸𝑟 (4.13)
In the above equations with reference to Figure 4.9, Rij and Wij are the pump and
signal stimulated rates, and the Aij are the nonradiative rates from level i to j. A21 is the
fluorescence rate. Cup is the homogeneous upconversion coefficient, and the C14 is
cross-relaxation coefficient. The Ni represents the ion population at each energy level,
and NEr is the total erbium ion population. The equations are normally solved
numerically using Runge-Kutta or similar methodologies.
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4.3.2 Simulation of erbium ion doped As2S3 waveguide amplifiers
Optiwave Corporation has commercially released a fully self-consistent model of Er3+-
Yb3+ codoped waveguide amplifiers based on the propagation and rate equations. This
is embodied in the Optisystem software suite and was employed here to implement the
simulation work. The amplification scheme used in the simulation is shown in Figure
4.10. A CW laser at 1538 nm with low output power of -30 dBm is used as a signal
source; a high power pump laser at 1480 nm with 200 mW output power is used as a
pump source; an optical spectrum analyser and WDM analyser are used to capture the
output spectrum and record the results. The core part of this design is the Er3+-Yb3+
codoped waveguide. Parameters required by simulation are measured from the
fabricated erbium ion doped As2S3 waveguides or from the literature [123, 257], and the
parameters used in the simulation are shown in Table 4.2.
Figure 4.10 Schematic optical set-up for Er3+-Yb3+ waveguide amplifier pumping
modelling using the OptiSystem software [258].
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Table 4.2 Parameters used in the OptiSystem software [258] simulation (some of them
cited from [123, 257]). Detailed definition of A32, A43, Cup and C14 could be found in Figure
4.9.
Name and description Value used in simulation
Waveguide length 3 cm
Signal background loss 1 dB/cm (@1300 nm)
Pumping background loss 1 dB/cm (@1300 nm)
Wavelength to calculate 1538 nm
Refractive index data file 2.41 @1550 nm
Er3+ ion density 6x1026 ions m-3
Yb3+ ion density 0 ions m-3
Er3+ metastable lifetime 1.3 ms
ESA cross-section value at 1480 nm 3x10-25 m2 [123]
A32 1x108 s-1
A43 2x104 s-1 [123]
Cup 4x10-23 m3 s-1 [257]
C14 2x10-24 m3 s-1 [257]
Noise centre frequency 1550 nm
Noise bandwidth 60 nm
Noise threshold -100 dB
Note: A32, the nonradiative rates from level 3 (upper level of 4I13/2 state) to level 2 (lower level of 4I13/2
state), is fixed as 1x108 s-1 to make sure the rapidly decay between these two levels.
As this model is designed for an Er3+-Yb3+ codoped waveguide, the Yb3+ density
was set to be 0 atoms/m3 in simulation for the absence of Yb3+ ion. It has to be pointed
out that in this model, any process occurring above the 4I9/2 state (known to occur here
as discussed in Section 3.4) is not taken into consideration, because short lifetime of the
4I9/2 state in silicate glasses. However, in chalcogenide glass due to the low phonon
energy, the lifetime of 4I9/2 state is long enough (intrinsic lifetime of 0.59 ms was
measured on an erbium ion doped As2S3 film, see Section 3.4) to support further up-
conversion processes, which could be confirmed from emission bands at 520 and 540
nm in the wideband emission spectrum described later in Section 4.4.2. Due to the
limitation of this model, influences of energy levels above the 4I9/2 state are not
considered in the following simulations.
Firstly, the small signal gain at three different signal wavelengths (1540 nm, 1555
nm and 1580 nm) were simulated for Er3+:As2S3 waveguides with different excitation
powers at 1480 nm. The simulation results are presented in Figure 4.11. Clearly, the
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results for the three different wavelengths share a similar trend, and the 1540 nm signal
has the best performance lying closest to the erbium ion gain peak. With a 1540 nm
signal, net gain can be achieved with pump power as low as 20 mW, and then the gain
saturates at its maximum value of 11dB at for a 100 mW pump. The 1555 nm signal has
a similar trend to the 1540 nm one, but less gain at 7 dB in total; while the 1580 nm
signal barely exhibits gain even with its best performance, which can be explained by
the relatively small emission cross-section of erbium ion in chalcogenide glass hosts at
this wavelength.
Figure 4.11 Simulated signal gain spectra versus pump power at different wavelengths of a
30 mm long Er3+:As2S3 waveguide amplifier with Er3+ concentration of 6x1026 ions m-3.
The gain spectrum under different excitation powers, ranging from 0.01 mW up to
200 mW, in the waveguide was also simulated and is shown in Figure 4.12. From the
results, the peak gain occurs at 1540 nm, and 50 mW can be considered the maximum
required pump power. Based on simulations, gain of 11 dB at 1540 nm in a 30 mm long
waveguide could be expected, corresponding to 3.67 dB/cm. The high level of
performance for low pump powers even with 1 dB/cm propagation loss is a
consequence of the high gain cross-section in chalcogenide glass hosts and the small
mode area afforded by the high refractive index contrast. Clearly performance of this
order would be attractive for compact devices such as amplifiers or laser arrays powered
off a single pump, or for ultrahigh repetition rate on chip mode locked lasers.
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Figure 4.12 Simulated gain spectra with different pump powers of an Er3+:As2S3
waveguide amplifier.
4.4 Waveguide fabrication
4.4.1 Waveguide fabrication introduction
At present, the usable length of chalcogenide planar waveguides is determined by the
loss of the waveguide at the pump wavelength. Reduction of the mode area, as required
to maximise nonlinear optics effects, is challenging since waveguide loss usually
increases super-linearly with the waveguide cross-section decrease, due to scattering
induced by side wall surface roughness [16, 259]. Besides careful design of the
waveguide structure to weaken the negative effects from rough surfaces and sidewalls,
another method to overcome this issue is improved process technology. Since the first
reports of waveguides made from chalcogenide glasses appeared in the 1970s from
researchers in Russia [260] and Japan [261, 262], various ways of fabricating
waveguides have been proposed. Some of the most popular methods are described
below.
(a) CW direct laser writing
Direct laser writing in chalcogenide materials due to their photo-darkening effect, which
refers to the phenomenon that when chalcogenide glass is illuminated by light of energy
above its bandgap, the optical absorption edge of illuminated material will shift to
longer wavelength. This property in chalcogenide glasses has been used to fabricate
waveguides since the 1970s. In 1979, Zembutsu et al. [262] reported waveguides based
on As-Se-S-Ge thin films using irradiation with a 1.06 µm Nd3+:YAG laser. The
refractive index change between the irradiated area and un-irradiated area was about
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0.03 at 1.06 µm wavelength which they believed was large enough to fabricate a 3-D
waveguide (curved waveguide and optical directional coupler). The best propagation
loss of 0.4 dB/cm at 1.064 µm was achieved in an As40Se10S40Ge10 (at%) film, and they
believed with an optimised film composition an even lower loss waveguide was
achievable. In Turnbull et al.’s work [263], 5 µm width waveguides were fabricated by
exposure to 514.5 nm CW light from an Ar3+ laser on sputtered Ge5As34Se61 (at%) films
with an approximate dosage of 100 J/cm2 on the sample. The propagation losses of the
obtained waveguides ranged 3.5 to 6.4 dB/cm at 1.3 µm, and scattering from micro-
cracks at the glass–substrate interface was speculated as the dominant source of the loss.
Nicolas et al. [264] reported a refractive index change of 0.04 induced in an As2Se3 film
by exposure to 633 nm He–Ne laser through a photomask. In waveguide fabrication, an
incident CW power of 1.0 mW, focused using an f=11 mm aspheric lens, at a translation
speed of 4 mm/min were thought ideal writing parameters, with which good lateral
confinement, single-mode output, and high injection efficiency were achievable. The
measured average propagation loss of the waveguides was 0.5±0.1 dB/cm for the TE
mode and 1.1±0.1 dB/cm for the TM mode at 8.4 µm, respectively.
Clearly, using the photodarkening property of chalcogenide glasses for waveguide
fabrication is a simple approach. However, with this method only low refractive index
contrast is achievable, and often the waveguides fabricated with this method are not
stable when heated or exposed to even long term ambient light.
(b) Femtosecond direct writing
Femtosecond lasers irradiating within the Urbach-tail region were also used to direct
write chalcogenide waveguides [265-267]. Zoubir et al. [207] wrote waveguides in
As2S3 films with femtosecond laser pulses from a 25-MHz repetition rate Ti:sapphire
laser, with an 40 nm spectral bandwidth centred at 800 nm, producing 20-nJ pulses of
30-fs duration with 4.5% pulse-to-pulse stability. In the waveguide, a refractive index
difference bigger than 10-2 was achieved. The structure change induced by Ti:sapphire
laser illumination was investigated using Raman spectroscopy. A buried single mode
waveguide within a Tm3+-doped 75GeS215Ga2S310CsI (mol% GGSI) glass substrate
was realised using femtosecond laser direct writing [267]. Single mode guidance at
1039 nm, with a minimum propagation loss of 0.86 dB/cm at 1039 nm, was reported. It
was pointed out that high stress could be induced in waveguides directly written using
low-repetition rate ultrafast laser systems [268].
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(c) UV direct writing
UV direct writing is also thought a useful way for waveguide fabrication. In Mairaj et
al.’s work [269], waveguides were written on Ga:La:S bulk glass with continuous wave
UV (244 nm) light providing spatially uniform fluence (±5%) in the range of 1.1 J/cm2
to 3.6 J/cm2. Propagation losses of the obtained waveguides was 0.2±0.1 dB/cm at 1.3
µm, with a positive change of refractive index of 10-3. Lin et al. [270] reported channel
waveguides with ~ 10-2 refractive index difference achieved by irradiating As2S8 films
under a UV light (300–436 nm) with an intensity of 58 mW/cm2 . After proper
annealing, propagation loss of 0.76 dB/cm was achieved at 1310 nm.
Similar to the photo-darkening and femtosecond methods, the refractive index
contrast obtained using UV direct writing is limited.
(d) Ion implantation
Ion implantation has also been trialled as a method for fabricating chalcogenide glass
waveguides. Meneghini et al. [271] fabricated channel waveguides on As2S3 films
through implanting helium ions at an energy of 113 keV, a dose of 2x1016 ion/cm2, and
an average current density of 1 µA/cm2. The ion range distribution calculated using
commercial software was 671 nm and the straggling width was 332 nm (half-width at
half-maximum). Also a metallic mask was applied in front of the film in order to define
the structure of the final waveguides. Confinement of 1.3 µm light in the implanted
channel was achieved, and this was explained by the refractive index increase of As2S3
due to helium implantation. However, the propagation loss of this waveguide was not
measured due to the too low level of the scattering loss over the buried channel
waveguide.
In Qiu et al.’s work [272], gallium lanthanum sulphide (GLS) and gallium
lanthanum oxysulfide (GLSO) glass waveguides were fabricated using Ar+ ion
implantation at 60 MeV and 2x1012 ions/cm2. The beam current density in the
experiment was chosen at about 10 nA/cm2 to minimise charging and heating effects
during irradiation. Waveguides with a refractive index distribution of ‘well + barrier’
type were obtained, where the ‘well’ region had an increased refractive index because of
the ionisation effect of electronic energy deposition, followed by the appearance of a
‘barrier’ region due to the energy deposition into nuclear displacements. Propagation
losses of 2.0 dB/cm for GLS and 2.2 dB/cm for the GLSO glass were achieved.
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It is well known that, ion implantation is an energy and time-consuming method,
and, with this method, the metallic mask used to define the waveguide structure is not
trivial to fabricate. Also the roughness caused by the high energy ion injected into
chalcogenide films is another issue needing careful consideration.
(e) Wet etching
Wet-etching is a well-known technology for waveguide fabrication from many host
materials, and it has also been applied in processing chalcogenide waveguides. In
DeCorby et al.’s report [273], waveguides based on As2S3 films were patterned directly
using photo-exposure through a mask followed by selective wet etching. Shallow ridge
waveguides (with ridge width of 3.8 µm), with losses as low as 0.26 dB/cm at 1530 nm,
were produced. As2S3 glass ridge waveguides were also fabricated with wet-etching in
Su et al.’s work [208]. The as-deposited As2S3 film was irradiated with a 2 mW 532 nm
laser with a spot diameter on the order of 2 µm. This 532 nm laser irradiation induced a
change not only in structure but in physicochemical properties as well, offering the
possibility that, with a properly chosen selective etching liquid, the etching rate for
exposed and unexposed parts could be very different. The film was then wet etched
using CS2 solution with iodine as an oxidising reactant. The average loss obtained for
the ridge waveguide was 5.5±0.5 dB/cm at 1330 nm. In this fabrication process no mask
was required.
However, the drawback illustrated by wet etching methods is that many
chalcogenide glasses are attacked by alkaline resist developers which leads to a
relatively rough sidewalls which in turn increases the scattering losses [250]. Precise
control of waveguide dimensions is another difficulty that needs to be overcome, along
with the non-ideal cross-sectional profile resulting from isotropic etching.
(f) Lift-off techniques
The lift-off method was proposed to overcome the issues caused by etch induced
sidewall roughness, in which a chalcogenide glass film is deposited onto a pre-patterned
resist structure with undercut. In Frantz et al.’s report [274], waveguides were patterned
by a lift-off technique on a gallium lanthanum sulphide (GLS) glass film deposited
using radio-frequency (RF) magnetron sputtering. The root mean square (RMS) surface
roughness of a slab waveguide was sufficiently small to provide low scattering losses.
Estimated propagation loss for the final lift off waveguide was 2.4±0.4 dB/cm at 1650
nm. High-index contrast Ge23Sb7S70 (at%) strip waveguides and small core Ge23Sb7S70
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(at%) rib waveguides (1.2 µm in width, 500 nm in height) have been fabricated using
the lift-off method in Hu et al.’s report [275]. The obtained strip waveguides had good
wafer-scale uniformity with loss of 2-6 dB/cm, while the rib waveguides had a lower
loss of 0.5 dB/cm.
Other processes, for example, thermal reflow, have been utilised to smooth the
sidewall roughness and then decrease loss. Although, with lift-off methods, the quality
and loss of waveguides are both improved, scattering from rough sidewalls is still
thought to be the main source of the loss.
(g) Reactive ion etching
Reactive ion etching refers to a method that removes material by exposing the materials
to a bombardment of ions which both chemically and physically remove material from
the exposed surface. In reactive-ion etching, high-energy ions are generated in low
pressure plasma, typically using an RF electromagnetic field. The ions bombard the
wafer surface and react with it forming volatile etch products that are pumped away.
Several etching parameters, such as pressure, gas flow, gas composition, and RF power,
can be chosen to meet different requirements and vary the degree of etch anisotropy.
High quality chalcogenide waveguides with low loss were fabricated using this method.
Ridge As2S3 waveguides produced using dry etching with CF4 and O2 have been
reported by Ruan et al. [12]. High index contrast rib waveguides (Δn~1) were obtained
with a minimum propagation loss of 0.25 dB/cm at 1.5 µm. Madden et al. subsequently
reported on the fabrication of etched As2S3 chalcogenide planar rib waveguides using
CHF3 and O2 [13]. Optical loss as low as 0.05 dB/cm at 1550 nm was achieved, which
is the lowest loss in chalcogenide waveguides ever reported so far and comparable to
losses obtained in even mature low index contrast germanosilicate technologies suitable
for dense integration.
Besides all the methods mentioned above, there are some other approaches that
could be used to fabricate chalcogenide glass waveguides, such as ion exchange [276];
fibre-on-glass [277]; physical sputter etching [278]; hot embossing [279, 280], etc.
However, each of them has its own drawback. For example, the fibre-on-glass and hot
embossing always need to heat the materials up to Tg, which will may enable the rare-
earth ions to cluster; with ion exchange an ionic species is required for exchange that is
not normally present in chalcogenides, and with physical sputter etching methods, it is
hard to get low loss chalcogenide waveguides preferable for amplification devices.
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4.4.2 Erbium ion doped As2S3 waveguide fabrication
Low passive losses in rare-earth ion doped planar waveguides are highly desirable as
they both maximise the external gain through lower signal losses, and also allow for
longer devices through lower pump passive losses with correspondingly higher
amplifier gain and saturated output power. Therefore, a method that could fabricate low
loss waveguides is important. To date, the best As2S3 waveguide with loss as low as
0.05 dB/cm was fabricated by UV lithography combined with inductively coupled
plasma reactive ion etching (ICP-RIE) [13]. The high quality waveguides and the
availability of facilities such as UV lithography and plasma etching made this method
attractive.
Based on the waveguide design discussed above (see Section 4.2.1), As2S3 bilayer
films were deposited by thermal evaporation on a <100> oriented 100 mm silicon wafer
with 2 µm of thermal silicon dioxide (TOX) as a bottom cladding: first an erbium ion
doped As2S3 layer with 800 nm thickness (erbium metal was used as the erbium ion
source), followed by an un-doped As2S3 layer of 550 nm. Before going to the UV
lithography, the as-deposited film was thermal treated at 130 °C for 24 hours followed
by green light treatment for 48 hours to push the film’s properties back toward its bulk
counterpart (see Section 3.2.3) [233, 234]. A 100 nm layer of SU-8 was spun and cured
on the film before putting it into the annealing oven. This SU-8 layer benefits films in
two ways: firstly, this thin SU-8 layer protects the films from O2 to prevent film quality
degradation by oxidation and surface diffusion [232]. Secondly, it is well known that
chalcogenide glass films are readily attacked by alkaline developers. Fortunately, this
SU-8 layer acting as a protective coating beneath the resist prior to patterning can
protect the films from this attack.
After thermal and green light treatment (see Section 3.2.3), the refractive index of
erbium ion doped As2S3 films can reach 2.41 at 1550 nm which is close to value of its
bulk counterpart (2.43). Nominal waveguide widths on the chrome mask were 1, 2, 3
and 4±0.1 μm. Photoresist coating and development was carried out in an SVG 8600
series track to ensure repeatability, and resist exposure was undertaken in a Karl Suss
Mask Aligner (MA-6) with a 350 W mercury arc lamp filtered to provide just the 365
nm i-line emission (~4 mW/cm2 intensity). The resist was 1 μm thick Clariant AZ MIR
701 and it was developed for 60 seconds in a single puddle using AZ 326 MIF
developer after a 1 minute post exposure bake at 110°C on a vacuum contact hotplate.
94
Following development, a 1 minute 110°C hard bake was applied in ambient
atmosphere.
All the etching work was performed in an ICP system (PlasmaLab 100 from Oxford
Instruments). This ICP system has a load lock chamber and sample chamber which are
pumped separately. The base pressure of the system is around ~2x10-6 Torr (~2.67x10-4
Pa) and the system can pump up to 100 sccm of gas and still maintain 8 mTorr (~1.1 Pa)
pressure. Gases used in etching were controlled by mass flow controllers. Etching depth
and etching rate were monitored live by an in-situ laser interferometer operating at 677
nm. The coated SU-8 layer was removed by exposing the film to O2 and Ar plasma in
the ICP system (200 W ICP power, 20 W forward power and 10 mTorr (~1.3 Pa)
pressure).
Previously, a combination of CF4 and O2 were employed to etch As2S3 [231].
However, the etching results indicated that this combination had different etching rates
for different phases present in the As2S3 films which is known to be inherently
nanoscale phase-separated [250, 281]. This leads to rough etched surfaces which are
undesirable. Based on experimental results, the main cause for this roughness was due
to the high proportion of reactive fluorine in CF4 which gave different etching rates for
some phases of the thermally post-treatment As2S3 host material. Therefore, CHF3 was
employed instead to reduce the concentration of fluorine and also to provide some
sidewall and surface passivation via the deposition of fluorocarbon polymers [13]. This
resulted in significantly lower waveguide losses. Thus, in this work, CHF3 was used as
the etching gas, and the details of the etching process was described in [13].
After etching, the residual resist was removed by using AZ Quikstrip wet stripping.
Normally, inorganic polymer is thought a good option for top protective cladding.
However, broad absorption bands inorganic polymer possesses are close to the
wavelengths of our interest, therefore, the waveguides were used with an air top-
cladding instead (see Figure 4.7). Whilst this made them more prone to damage and
limited their life, it was deemed an acceptable compromise in most testing instances.
After hand cleaving with a diamond scriber, the waveguides were ready to measure. A
schematic of the whole fabrication procedure is shown in Figure 4.13. Figure 4.14
shows an optical micrograph of the end facet of a finished 2 µm width waveguide.
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Figure 4.13 As2S3/Er3+:As2S3 (after thermal post-treatment) waveguide fabrication
procedure using photolithorgraphy. TOX refers to thermally oxidized silicon.
Figure 4.14 Image of the end facet of a 2 µm width Er3+:As2S3 waveguide depicted in
Figure 4.7.
4.5 Characterization of erbium ion doped As2S3 waveguides
To test the waveguides both passively and in a pumped configuration, an in-house
supercontinuum (SC) source was used as a probe/signal source. The SC source was
generated in a photonic crystal fibre driven by 10 ps pulses from a mode-locked
Nd3+:YVO4 laser [282, 283]. This was coupled into the waveguides through a lensed
fibre, with wavelength range from 300 nm to 2000 nm and total power up to around 200
mW. Transmission spectra using this SC source were recorded using an optical
spectrum analyser (Ando AQ 6317) and the supercontinuum spectrum is shown in
Figure 4.15. The SC source was heavily attenuated (<-30 dBm/nm) during the
experiment to avoid bleaching the Er3+ absorption in waveguide samples.
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Figure 4.15 Spectrum of the SC source used as signal source in this thesis.
4.5.1 Propagation loss and absorption of erbium ion doped As2S3 waveguides
The first concern regarding the obtained Er3+ doped As2S3 waveguide is its propagation
loss and Er3+ absorption spectrum, which can both be derived from loss spectra of the
waveguide. Loss spectra were measured using the SC system described above (see
Section 4.5) as a probe source, and recorded with the described OSA (see Section 4.5)
on the output side of the waveguides. The set-up for this measurement is shown in
Figure 4.16. The attenuated SC source was delivered to waveguides by a lensed fibre
from Oz Optics. The lens focused the SC source to a spot size of 2.5 µm (1/e2) diameter
at 14 μm distance from the tip. On the output side of the waveguide, another lensed
fibre was positioned to collect the output signal. Both lensed fibres were mounted on 3
axis piezo enhanced micro-positioning stages for hands-off alignment. Reference fibre
to fibre SC spectra were recorded at the start and end of each measurement session to
mitigate any drift in the SC spectrum shape.
Figure 4.16 Optical set-up for Er3+:As2S3 waveguides loss measurement.
97
The typical insertion loss spectrum of a 180 mm long, 3 µm width serpentine
waveguide, with 2 mm radius bends, with 0.15 at% erbium ion is shown in Figure 4.17
after normalising out the SC spectrum using a lensed fibre to lensed fibre measurement
as the reference. Two absorption bands are visible in the spectrum; one located at 980
nm, corresponding to the absorption of 4I15/2→4I11/2 transition; and the other located at
1540 nm which is ascribed to the absorption of 4I15/2→4I13/2 transition.
Figure 4.17 The insertion loss spectrum of a 180 mm long 3 µm width Er3+:As2S3
serpentine waveguide with 0.15 at% erbium ion concentration. The red fitted curve
represents the loss due to Rayleigh scattering.
Besides the erbium ion absorption bands, the background also provides information
on the loss of this waveguide at different wavelengths. The waveguide loss spectrum
indicates an estimated background propagation loss of about 0.35 dB/cm at 1550 nm
after correcting for coupling losses (calculated at 2.35 dB/facet from mode overlap
using a Matlab code based on the finite difference method and reflection [253]), and
shows an excellent fit to a 1/λ4 curve, where λ is the wavelength. This indicates that the
loss is dominated by Rayleigh scattering off nanoscale inhomogeneities and is unusual
for As2S3 waveguides which typically show loss curves dominated by 1/λ2 responses
from sidewall scattering. From the fit, the erbium ion absorption was extracted. The
absorption band at 1540 nm is 34 dB in this 180 mm long waveguide, leading to a 1.9
dB/cm absorption of erbium ion in this 0.15 at% (0.45x1020 ions/cm3) doped
waveguide.
The absorption band centred at 1540 nm was also extracted, normalized and is
plotted in Figure 4.18. A normalized erbium ion absorption band for a TeO2 waveguide
[152] is also plotted in Figure 4.18 for comparison. One immediately striking difference
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is in the short wavelength absorption edge, where with As2S3 as a host, the absorption at
the usual erbium ion pump wavelengths in the 1480-1490 nm range is considerably
weaker than in TeO2. This suggests that a pump wavelength in the 1505-1510 nm range
may be required for efficient pumping. The wavelength of the maximum absorption is
also red shifted as expected, and this could be explained by the nephel-auxetic effect
[39], which was discussed in Section 3.3.2(c).
Figure 4.18 Normalized absorption bands from 1450 to 1650 nm in Er3+:As2S3 and
Er3+:TeO2 waveguides. The black vertical arrow shows the position of the usual Er3+ 1480
nm pumping wavelength.
4.5.2 Amplification measurements of erbium ion doped As2S3 waveguides
In the interests of clarity, it is noted that the raw internal gain of a waveguide amplifier
is considered to be the enhancement factor minus the rare-earth ion absorption, where
the enhancement factor is the ratio of the output power with pump on, minus the
amplified spontaneous emission (ASE) to the output power with the pump off. This is
the definition used later to measure the internal gain and represents actual signal
amplification due to the material processes down the waveguide length indicating the
likely gain available if propagation losses are low (presumed to be the case for any
useful amplifier). Raw transparency occurs when the internal gain equals 1, that is it
offsets the rare-earth ion absorption exactly, and gain exists when it exceeds unity. Net
internal gain results when the division by the waveguide propagation losses leaves a
number above unity and means more power exits the waveguide than entered it.
Excitation experiments on waveguides were performed with the set-up shown in
Figure 4.20. The excitation laser was centred at 1490 nm with a maximum output power
of 200 mW exiting the SMF-28 output connector. This was combined with the
attenuated SC source (see Section 4.5) which acted as a probe signal via the 10% port of
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a 10/90 coupler and then was delivered to the waveguide through a lensed fibre as
previously described. On the output side of the waveguide, the signal was collected by
another aligned lensed fibre, then was passed through two 1490/1550 nm WDMs which
provided roughly 30 dB attenuation for the excitation power at each WDM. After
removing the remnant excitation power, the signal spectrum was recorded by an OSA.
A bi-directional pumping setup can also easily be formed by adding another excitation
laser to the output side via the first WDM.
Figure 4.20 Optical set-up for excitation experiments on Er3+:As2S3 waveguides with 1490
nm uni- or bi-directional pump with a SC probe source (see Figure 4.15 for output
spectrum of this SC source).
It is worth noting that in the set-up described here, the loss from the lens tipped
fibres themselves is about 1.8 dB (which was zeroed off to a lens tipped fibre to lens
tipped fibre measurement, consequently it had no effect on the waveguide insertion
loss), and the overlap loss due to the mismatch of the modes between the lensed fibre
and the waveguide and surface reflection are about 2.35 dB/facet in total based on
calculation. Therefore, this set-up comes with a total insertion loss of about 6.5 dB.
The test results are plotted in Figure 4.21 for a 55 mm long, 2 µm width, 0.15 at%
Er3+ doped As2S3 straight waveguide uni-directionally pumped at 1490 nm. The sharp
drop of power below 1520 nm and above 1640 nm is due to the cut-off of the 1490/1550
nm WDM applied here. The erbium ion absorption band at 1540 nm bleaches with
increase of excitation power. From Figure 4.21, excitation of 70 mW at 1490 nm in the
waveguide gives the best performance of this waveguide. Any further increase of
pumping power only resulted in photosensitive, where the spectrum is degraded with
additional dips and spikes, and thus spectra with higher excitation power are not given
in Figure 4.21.
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Figure 4.21 Excitation results for a 55 mm length, 0.15 at% Er3+:As2S3 waveguide, uni-
directionally pumped at 1490 nm with SC probe (see Figure 4.15).
To reiterate, enhancement is the ratio of output power with pump on, minus the ASE,
to the output power with the pump off. Therefore, to get a clear enhancement, ASE
spectra also have to be recorded under the identical condition only with the SC source
turned off. The same set-up was employed to measure three spontaneous emission
spectra under excitation of 9 mW, 16 mW and 70 mW, respectively, and the results are
shown in Figure 4.22.
Figure 4.22 Amplified spontaneous emission (ASE) spectra for a 55 mm length, 0.15 at%
erbium ion doped As2S3 waveguide, with uni-directional pump at 1490 nm.
Enhancement factor spectra were then extracted from the excitation spectra, ASE
spectra and the zero excitation power spectra. The erbium ion absorption curve was
obtained from the insertion loss spectrum and the waveguide propagation loss fitting
curve. This was then applied to the enhancement factor curve to obtain an internal gain
curve, which is shown in Figure 4.23.
101
Figure 4.23 Internal gain characteristics for different 1490 nm pump intensities working
uni-directionally in a 0.15 at% Er3+:As2S3 waveguide of 55 mm length.
A small internal gain (~0.7 dB) from 1570 to 1630 nm was observed, the first time
to the Author’s knowledge that any internal gain has ever been reported in an erbium
ion doped chalcogenide bulk glass, fibre or thin film based waveguide amplifier. That
internal gain shows first at the longer wavelengths can be explained simply by the
emission cross-section inhomogeneous broadening and the correspondingly low
absorption cross-section there. However, it is clear that at the absorption peak
wavelength 1538 nm, the device is still about 4.3 dB below the optical transparency line
even for the best performance spectrum in Figure 4.23. Considering the 10 dB
absorption at 1538 nm of this 55 mm long waveguide, it is reasonable to estimate that
roughly 30% of total erbium ions are inverted in this waveguide. Details of the erbium
ion inversion estimation method can be found in Section 6.4.
It was not possible to utilize the full pump power as photosensitive effects began to
occur, characterized by the presence of mode coupling gratings and the appearance of
deep mode coupling dips in the spectrum. Whilst this complicated matters, it was
however clear that the gain spectrum was not increasing at any significant rate with
additional pump power. Similarly switching to 1505 nm pumps did not produce a better
result.
For comparison, an 18 mm length, 0.6 at% Er3+:As2S3 waveguide was also pumped
under the same conditions and the result is shown in Figure 4.24. About 8 dB optical
enhancement is found at 1540 nm when only 5 mW pump power applied, which is
believed to have been due to the ASE. No further optical enhancement occurs even
when the pump power is pushed up to 100 mW. Also it is clear the ASE level of this
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sample (~-65 dBm) is much lower than that of the 0.15 at% Er3+:As2S3 sample (~-
51dBm), and both of these values were measured with 2 nm optical bandwidth.
Figure 4.24 Excitation results for an 18 mm length 0.6 at% Er3+:As2S3 waveguide with uni-
directional pump at 1490 nm.
Given the evident non bleachable absorption in the excitation spectra, it is suggested
that there could be significant clustering in this 0.6 at% Er3+:As2S3 waveguide. Even the
0.15 at% Er3+:As2S3 waveguide appears to have a significant degree of clustering and
further improvements to the deposition process are required to attain net gain.
4.6 Co-thermal evaporation of neodymium ion doped As2S3
It is well known that it is more difficult to achieve optical gain in a quasi-three-level
energy system like Er3+ than in a four-level energy system like Nd3+, due to the absence
of the ground state absorption at the emission wavelength in a four level energy system.
Referring to Figure 4.25, Nd3+ ions in the 4I9/2 ground state absorb 808 nm photons and
are excited to the 4F5/2 state, and then by phonon emission rapidly decay to the upper
lasing 4F3/2 state. 1.06 µm emission occurs when ions transfer from the 4F3/2→4I11/2
levels, and the ions in the 4I11/2 level are quickly de-excited to the ground state 4I9/2,
again, due to a phonon mediated process. Due to the long lifetime of the 4F3/2 state and
the efficient nonradiative decay between the 4I11/2 and the 4I9/2 energy levels, a
population inversion is easily formed between the 4F3/2 and the 4I11/2 energy levels. High
gain is possible due to the lack of absorption at the emission wavelength. Therefore, it
would be interesting to know whether population inversion in Nd3+ doped As2S3
waveguides could be obtained.
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Figure 4.25 Energy level diagram of Nd3+ ions.
4.6.1 Neodymium ion doped As2S3 film deposited using co-thermal evaporation
Much work on neodymium doped chalcogenide glasses, fibres and waveguides has been
reported. The first attempt of measuring absorption and emission spectra of neodymium
ions in chalcogenide glasses was made by Reisfeld et al. in 1977 [42]. Chalcogenide
glasses of composition 0.7Ga2S3-0.27La2S3-0.03Nd2S3 (mol%) were prepared.
Absorption and emission spectra of these glasses were measured and compared to
commercial neodymium ion doped silicate glass. Higher absorption intensity and lower
nonradiative relaxation rate were observed compared with their silicate counterparts
[42]. In Mori et al.’s paper [284], Nd3+ doped Ge–As–Ga–Sb–S (core) and Ge–As–S
(cladding) fibre was drawn by the rod-in-tube method. Under 890 nm pumping, internal
gain of 6.8 dB at 1.083 μm was achieved from a 50 mm length fibre with doping
concentration of 1000 ppm. Net signal gain was not obtained due to the high
transmission loss of the fibre. Laser action in a neodymium ion doped chalcogenide
glass was demonstrated for the first time in a Ga-La-S host glass by Schweizer et al.
[105]. 1.083 μm laser action was achieved with a 1.42 mm thick sample of Ga-La-S
glass doped with 1.5 mol% Nd2S3 (2.6xl020 ion/cm3) neodymium ions, under excitation
wavelengths of 815 nm and 890 nm from a Ti:sapphire laser. In the following year,
laser action at 1.08 μm from a 22 mm long 0.05 mol% Nd2S3 doped Ga-La-S fibre
fabricated using the rod-in-tube technique was reported by the same group [106]. In
Mairaj et al.’s report [108], channel waveguide structures were directly written using a
focused UV-laser beam in neodymium ion doped Ga-La-S glass. A lasing experiment
was performed on a 0.5 mol% Nd2S3 (8.7xl019 ion/cm3) doped 16 mm long waveguide.
104
Single mode laser operation at 1075 nm with output slope efficiency of 17% with
respect to absorbed power was obtained. Attenuation in the laser device was estimated
by the Findlay–Clay method combined with a separate calculation involving the slope
efficiency, and the upper limit of the propagation loss was found to be about 0.5 dB/cm.
Most of the work has focused on the Ga-La-S glass system because of its high rare-earth
ion solubility, and there were few reports on other more frequently used glasses for
waveguides such as As2S3.
In this work, Nd3+ doped As2S3 films were deposited in a similar manner to that of
the co-thermally evaporated erbium ion doped As2S3 films, with the substitution of
neodymium powder instead of erbium metal. Evaporation was performed at a vacuum
level around ~3x10-7 Torr (~4x10-5 Pa). After reaching the pressure set point, a slow
increase of the neodymium powder temperature showed evaporation commenced at a
temperature of ~ 1000 °C, and the rate reached 0.001 nm/s at 1130 °C. For the film
deposition, the neodymium temperature was set at ~ 1060 °C and the rate of As2S3 was
set at 0.1 nm/s to target a final neodymium ion concentration of ~0.3 at%.
After deposition, a 120 nm SU-8 protective layer was coated on the obtained
Nd3+:As2S3 film, then thermal post-treatment (at 130 °C for 24 hours under vacuum)
and light post-treatment (500-570 nm green light with ~3mW/cm2 intensity for 48
hours)were carried out. The obtained film was characterized using the SCI Filmtek
4000. After thermal and light post-treatment, the film refractive index was 2.41 at 1550
nm and the bandgap was 2.27 eV, measured from a 1.48 μm thick film. These results
indicated the Nd3+ doping had little effect on the host glass As2S3 which exhibited a
refractive index of 2.41 at 1550 nm and bandgap of 2.34 eV (see Section 3.23).
4.6.2 Characterization of the neodymium ion doped As2S3 film
Based on the energy level diagram of Nd3+ ions shown in Figure 4.25, the lifetime of the
upper 4F3/2 energy level related with the 1.08 μm emission is of most concern. However,
due to lack of high isolation pump/signal WDMs working at this wavelength, the 4F3/2
energy level lifetime could not be measured.
The Nd3+:As2S3 film emission spectrum was collected using a Horiba Jobin Yvon
64000 spectrometer. The excitation laser operated at 830 nm, which is not the optimum
wavelength to excite Nd3+ ions from the ground state up to the 4F5/2-2H9/2 levels, but it
still lays in the long wavelength edge of the absorption band. An InGaAs detector
105
cooled using liquid nitrogen was installed on the spectrometer to record the emission
spectrum. The PL spectrum from a 1.3 μm thick neodymium ion doped As2S3 film is
recorded in Figure 4.26.
Figure 4.26 Emission spectrum of a 1.3 μm thick Nd3+ doped As2S3 film pumped at 830 nm.
There are two main features in the emission spectrum of the Nd3+:As2S3 film: the
major one is located at 1080 nm with broad shoulder arising from the 4F3/2 → 4I11/2
transition, and another band centred at 916 nm is associated with energy decay from the
4F3/2 to the 4I9/2 levels (see Figure 4. 25). The distinct broadening of the fluorescence
band of the long wavelength 1080 nm emission is thought to be due to the variation of
environments surrounding the rare-earth ions [285, 286], that is non-homogeneous
broadening. The same phenomenon is also noticed in silica-based glass. For example, in
Ainslie et al.’s report [287], significant emission spectrum broadening was observed in
some of the high Nd3+ doped (up to 15 wt% of NdCl3) SiO2-A12O3-P2O5 glasses. Also,
the peak wavelength at 1.08 μm, is roughly 10–20 nm longer than that found for Nd3+-
doped silicate, fluoride and phosphate-based glass fibres [284]. A similar wavelength
red shift is also noticed in erbium ion doped chalcogenide glass hosts [238], which is
discussed in Section 3.3.2(c).
Similar to the approach used to pattern the erbium ion doped As2S3 films to make
waveguides, two layers of neodymium ion doped and un-doped As2S3, respectively, of
designed thickness were deposited on thermally oxidized silicon (TOX) wafers,
followed by standard lithography and plasma etching with CHF3. Ridge waveguides
with the structure shown in Figure 4.27 were achieved.
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Figure 4.27 Typical structure of a 2 µm Nd3+ doped As2S3 waveguide (Tox is thermally
oxidized silicon).
Propagation losses and the neodymium ion absorption spectrum of the obtained
waveguides were measured using the SC source described in Section 4.5. To avoid
bleaching the neodymium ion absorption, a 20 dB optical attenuator was connected in
the beam path before the SC light was launched into the waveguides. The transmission
spectrum of this SC source was then recorded with, and without, the waveguide present
using an optical spectrum analyser (Ando AQ 6317) at the waveguide output side. The
transmission spectrum of an 8 mm long, 3 μm wide waveguide is shown in Figure 4.28.
Figure 4.28 Absorption spectrum of an 8 mm long, 3 µm wide Nd3+ doped As2S3 waveguide
after thermal and light post-treatment.
Two absorption bands are clearly observed in the absorption spectrum of
Nd3+:As2S3 waveguide. The strongest one is located at ~820 nm corresponding to the
absorption of ions from the 4I9/2 to the 2H9/2-4F5/2 excited levels. The band centred at 890
nm arises from the transition 4I9/2→4F3/2 (see Figure 4.25). More absorption bands at
shorter wavelengths, associated with exciting the ions to even higher energy levels, are
not observed in this spectrum because of the intrinsic bandgap absorption in the host
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glass. The absorption band at ~820 nm, with a broad shoulder (from 790 nm to 850 nm),
enables the PL spectrum to be obtained with excitation of 830 nm as mentioned before
(see Section 4.6.2). The ~820 nm absorption band is so strong that only a few photons
were collected by the detector making the spectrum noisy, therefore, it is not possible to
quantify its absorption strength. From Figure 4.28, the 890 nm absorption band is ~ 32
dB/cm in this 8mm long Nd3+:As2S3 waveguide. Comparing this value with the
absorption information in ref [105], the estimated Nd3+ concentration in this As2S3
waveguide is ~ 0.4 at% (1.2x1020 ions/cm3), which is a little higher than the number
~0.3 at% (0.9x1020 ions/cm3) estimated from film deposition rate. With such a high
absorption at the pump wavelength, the effective pumped length is short, which brought
practical difficulties in making accurate gain measurements.
The propagation loss of this waveguide could also be estimated from the insertion
loss curve. The total insertion loss for this 8 mm long waveguides is 4.9 dB at 1300 nm,
a wavelength far away from the absorption bands. Subtracting the fibre-to-waveguide
coupling loss of 1.55 dB/facet due to mode mismatch calculated by the finite difference
method, and a reflection loss of 0.75 dB/facet from the waveguide-to-air interface, the
propagation loss of this waveguide is estimated to be about 0.4 dB/cm. Although it is
not realistic to calculate waveguide propagation loss accurately with such a short
waveguide, due to uncoupled light bleed through, from the result it is still safe to say the
propagation loss of this waveguide is reasonably low and suitable for a potential
waveguide amplifier/laser. In the range from 1 μm to 1.7 μm, the propagation loss
shows almost no dependence on the wavelength. This phenomenon is different from
that in the un-doped As2S3 waveguides, where 1/λ2 dependence is typical indicating the
effects from roughness of the waveguide sidewalls [280]. It is also different from the
trend in the previous erbium ion doped As2S3 waveguide, where a strong 1/λ4
relationship is evident indicating nanoscale scattering off film inhomogeneities (see
Section 4.5.1). The flat curve obtained here may indicate a high quality waveguide with
smooth sidewalls and of a homogeneous nature at the nanoscale, or may arise due to
wavelength independent scattering from an unknown source.
Having no suitable WDMs, gain measurements were not feasible, so a lasing
measurement was attempted using an 806 nm diode laser pumping an 8 mm long 2 μm
wide waveguide (structure of this waveguide is depicted in Figure 4.27) with an optical
set-up as in Figure 4.29. The 806 nm laser with a beam diameter of 2 mm was coupled
into a single mode fibre (SMF 28) via a 10x objective lens with NA of 0.2. Then the
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excitation laser was delivered to the waveguides through a lensed fibre. On the output
side, another lensed fibre was positioned to collect all the signals from the waveguide,
and an OSA was connected to this lensed fibre to record all the data. The laser cavity
was formed by the facet reflections of the waveguide.
Figure 4.29 Optical set-up for lasing measurement in a ~0.4 at% Nd3+ doped As2S3
waveguide with 806 nm pump (OSA is optical spectrum analyser).
The obtained ASE spectra under different excitation powers are shown in Figure
4.30. The sharp peak at 808 nm is the residual pumping energy, and the bands at 922 nm
and 1080 nm correspond to the emissions of 4F3/2→ 4I9/2 and 4F3/2 → 4I11/2, respectively.
Figure 4.30 Amplified spontaneous emission (ASE) spectra of an 8 mm long, 2 μm wide,
~0.4 at% Nd3+ doped As2S3 waveguide with 808 nm excitation at increasing power.
Unfortunately, except for the ASE spectrum (see Figure 4.30), no lasing was
observed even under the maximum excitation power. The reasons for this less
preferable performance may be multi-fold. Firstly, the reflectivity of each end facet of
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the waveguide is estimated to be 15.7 % (-8.04 dB), which introduces significant
intracavity loss for a waveguide laser. However, given the doping level (0.4 at%) of
Nd3+, sufficient gain should have been available to overcome this. Secondly, the high
pump absorption (which could not be properly quantified) leads to a short effective gain
length, but the low output pump power indicating no bleaching in excitation which is a
signature of clustering. Lastly, from the spectrum, the absolute ASE level of this
waveguide is quite low even compared to erbium ion in an As2S3 film (compare ~-50
dBm for Er3+:As2S3 (Figure 4.22) with ~-60 dBm for Nd3+:As2S3 (Figure 4.30) with 2
nm optical bandwidth), although due to the lack of ground state absorption meaning it
should be higher ASE level than that of Er3+:As2S3. Considering the low rare-earth ion
solubility in As2S3, and the broadening of the PL emission shoulder from the film PL
spectrum, neodymium ion clustering is possible in this case.
Given the excitation results of Er3+ & Nd3+ doped As2S3 waveguides, it is highly
suspected that rare-earth ion clustering occurs during the film deposition. The
evaporation of metals (more than 40 elements by 1969) has been studied in detail in the
past [288], and it is established that during evaporation from the metal, single atoms are
emitted together with polyatomic complexes, which become overgrown with atoms in
the region of high vapour density close to the source. In addition, as the plume of
evaporant expands and cools, polyatomic complexes could be formed by atomic
collisions if conditions are favourable to the removal of the heat of condensation.
Although no direct evidence was found to prove the existence of polyatomic rare-earth
complexes during film deposition, the non bleachable Er3+ absorption in 0.6 at% Er3+
doped As2Se3 waveguide, and small rare-earth particles formed on the edge of alumina
crucible after evaporation, indicating the high possibility of polyatomic evaporation.
Based on the research [289], the relative equilibrium vapour pressure of the dimer, P2, is
given by the expression:
𝑃2
𝑃= 𝑐𝑜𝑛𝑠𝑡𝑒
∆𝐻0𝑅𝑇
(1−𝛼
𝛼) (4.14)
where Ρ is the vapour pressure of the monomer, R is the gas constant, Τ is the absolute
temperature, and α=ΔH0/D0, where D0 is the dissociation energy of the dimer, and ΔH0
is the heat needed to evaporate an atom at 0 K. As for metals, α>1, means the proportion
of dimers in the saturated vapour increases with temperature. Therefore, working on this
premise, it is suggested that a lower evaporation temperature would: a) reduce the
percentage of evaporated polyatoms and condensed clusters through a less dense plume,
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and b) a lower rare-earth ion concentration would reduce the likelihood of any post-
treatment induced clustering. However, to achieve a reasonable rare-earth ion
concentration with lower evaporation temperature means a large area source is required,
certainly much larger than the half inch diameter crucible source used here. Another
option is to use a plasma assisted reactive evaporation device that combines an
evaporation source and a high density plasma source [290]. These two sources are
installed in a coaxial configuration, where the material evaporated from the evaporation
source is transported through the plasma source and then condenses on the substrate.
With the help of the plasma, the dimers can be dissociated to ensure monatomic rare-
earth ion deposition. However, limited by the chamber and facility available, it was not
possible to try these options. Consequently, no further progress was made on thermally
evaporated films.
4.7 Conclusion
In this chapter, the basic theory of planar optical waveguides was discussed. Methods
for waveguide mode simulation were introduced, and the amplification performance of
an erbium ion doped As2S3 waveguide amplifier was simulated by Optisystem based on
rate equations and the waveguide simulation data. Erbium ion doped As2S3 waveguides
with loss of 0.35 dB/cm at 1.55 µm were fabricated using standard photolithography
and plasma dry etching. With the obtained 55 mm length, 2 µm wide, 0.15 at%
Er3+:As2S3 waveguides, for the first time, internal gain from 1570 to 1630 nm was
achieved in an erbium ion doped chalcogenide glass or in a chalcogenide thin film based
waveguide amplifier. ASE signals at 900 nm and 1100 nm were observed in
neodymium ion doped As2S3 waveguide fabricated using thermal evaporation.
However, due to the lack of proper wavelength division multiplexing (WDMs) and
pump laser, lasing in this waveguide was not achieved. Erbium ion clustering and low
rare-earth ion solubility in thermal evaporated As2S3 film are thought the main reasons
for the limited internal gain achieved in the amplification experiment. Therefore,
different film fabrication approach (RF sputtering in Chapter 5) and different host
material with better rare-earth ion solubility (Ge-Ga-Se material in Chapter 6) are tried
in the following chapters.
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Chapter 5
Different approaches to fabricate erbium ion
doped As2S3 waveguides
Although limited internal gain was achieved from 1570 to 1630 nm in an erbium ion
doped As2S3 waveguide deposited using thermal evaporation with metallic erbium as
the dopant, the maximum erbium ion inversion was estimated at around 30%. Higher
erbium ion inversion is required to make practical devices. From excitation experiments
and high temperature thermal post-treatment results in Chapter 4, the low erbium ion
population inversion level appeared to be due to erbium ion clustering in As2S3 films.
Therefore, in this chapter, several different approaches are introduced to attempt to
eliminate the erbium ion clustering.
5.1 Co-thermal evaporation of ErXn (X=S/Cl) with As2S3
It is known that in rare-earth ion doped chalcogenide materials, films offer certain
opportunities over bulk glasses in hosting rare-earth ions. In bulk glass fabrication, the
melt-quenching method with long rocking time provides high quality homogeneous
glasses. At the same time, long rocking time also offers rare-earth ions time and energy
input to migrate. Therefore, the rare-earth ion solubility in a host material caps the limit
on the highest concentration of rare-earth ion that can be incorporated in the host glasses
without clustering. On the other hand, thin films (by thermal evaporation or other
methods) are usually produced in a non-equilibrium condition by rapidly condensing
vapor onto a cold substrate which gives insufficient time for the material to revert back
to their metastable equilibrium state. Therefore, higher concentrations of rare-earth ion
than the actual solubility limit could be incorporated, provided that the film is not then
subjected to elevated temperatures or other forms of energy input that can cause bond
rearrangement or material diffusion. Therefore, as discussed in Chapter 4, the clustering
in the low concentration of erbium ion doped sample (0.15 at%) stemmed from the
metal erbium evaporation source. One possible way to prevent the formation of erbium
metal clustering is to use erbium compounds instead of erbium metal in film
evaporation. By using compounds such as ErCl3 and Er2S3, the direct formation of
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erbium ion clustering should be reduced due to the erbium ion being already bonded
with -Cl or –S.
For this experiment, the two candidates that could be used to replace erbium metal
are ErCl3 and Er2S3. ErCl3 is the obvious first choice as it has a lower evaporation
temperature than Er2S3 (melting temperature of ErCl3 is 1529 °C compared to 1730 °C
for Er2S3). In the deposition chamber the evaporation temperature should be somewhat
below the standard melting temperature due to the low pressure. It was reasonable to
assume that for the same evaporation rate, ErCl3 would need a lower temperature than
Er2S3 does. This provides more space to operate and less potential heat load in the
chamber. The maximum operating temperature for the available alumina crucibles used
in the evaporation was 1500 °C. Thus, ErCl3 was the first to be deposited to test the
hardware capability.
5.1.1 Co-thermal evaporation of ErCl3 with As2S3
ErCl3 has been used as a source of Er3+ in various host materials. For example, Shojiya
et al. [291] studied the upconversion luminescence from Er3+ in chloride glasses using
ErCl3 as an active Er3+ source. Photoluminesence properties of Er3+ in oxyfluoride glass
ceramics were investigated by Zeng et al. [292], and ErCl3 was one of the erbium
compounds used in this report. Soundararajan et al. [293] reported the optical properties
of erbium ion doped fluorochlorozirconate glasses, where active Er3+ were added as
ErCl3.
There are only a few papers on ErCl3 utilization in chalcogenide glass hosts. In
Allen et al.’s work on the photoluminescence measurement of erbium ion doped GaGeS
and GeGaSe glasses [119], ErCl3 was employed instead of Er2S3 for its lower melting
temperature, thereby simplifying the glass preparation. However, the Er3+ introduced as
ErCl3 had less promising performance: no discernible lifetime was observed in the
ErCl3-doped samples, and the PL spectrum showed limited optical activity of Er3+ ion.
In reality, researchers are reluctant to use ErCl3 in chalcogenide glasses as the Cl- in the
host might act as a quenching centre and it is considered to cause instability in the glass
structure [119, 294]. Despite Allen’s results, it is decided to try using ErCl3 in this
instance as the experiments are with As2S3 not Ga-Ge-S/Se.
High purity (5N elements) As2S3 and ErCl3 components were loaded into a
resistively heated tantalum boat and resistively heated alumina crucible, respectively.
113
Evaporation was performed at a pressure ~1x10-5 Pa. After reaching the pressure set
point, the ErCl3 temperature was increased slowly; at the same time the rate sensor for
ErCl3 was carefully monitored to record the ErCl3 evaporation temperature. According
to the rate sensor (a quartz microbalance thickness monitor), ErCl3 started to evaporate
at about 620 °C. During deposition, the ErCl3 temperature was fixed at 650 °C to get a
stable evaporation rate of 0.001 nm/s, whilst keeping the As2S3 rate at 0.1 nm/s to
ensure the final Er3+ concentration was around 0.2 at% (~0.6x1020 ions/cm3). After all
the desired evaporation rates were achieved, the shutter covering the wafers was opened
and evaporation started. When the film thickness read from the film thickness monitor
reached the final thickness set point, the system closed the shutters for both sources and
films automatically, ramped down the power for both sources gradually and allowed
them to cool to <100°C before venting the chamber.
To protect the As2S3 from surface crystallisation, a 100 nm thick SU-8 layer was
coated on the surface and cured as soon as the as-deposited films were removed from
the chamber. Another set of films with a higher Er3+ concentration of ~0.4 at%
(~1.2x1020 ions/cm3) were deposited in the same way by doubling the ErCl3 evaporation
rate from 0.001 nm/s to 0.002 nm/s.
Optical properties of the films were measured using the FilmTek 4000, and the
lifetimes of the erbium ion doped As2S3 films were measured in the all-fibre confocal
setup described previously (in Section 3.3.2) under 1490 nm excitation. The results for
the as deposited film are listed in Table 5.1.
Table 5.1 Physical properties of as-deposited Er3+:As2S3 films with ErCl3 as dopant.
Er3+ concentration in
As2S3 film
Thickness /
nm
Bandgap / eV RI @ 1550 nm
Low doped (0.2 at%) 931 2.29 2.32
High doped (0.4 at%) 718 2.30 2.32
The film refractive index of 2.32 at 1550 nm and the bandgap around 2.3 eV are
identical to the values previously obtained from as-deposited As2S3 films (see Section
3.2.3), indicating that low concentration of ErCl3 has no clear effects in As2S3 film
formation.
1/e, intrinsic lifetime and PL intensity of the Er3+:As2S3 films under 20 kW/cm2
excitation were measured, and are shown in Table 5.2. From the results, 0.2 at% ErCl3
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doped As2S3 film has the best lifetimes, which are 2.5 ms (1/e) and 2.6 ms (intrinsic
lifetime), respectively. The lifetime drops with doubling of the Er3+ concentration, but is
still 1.4 ms, longer than the previous best data of the erbium metal doped 0.15 at%
As2S3 film. The much longer lifetimes obtained in ErCl3 doped films implies less
erbium ion clustering is formed in this ErCl3 doped As2S3 film than in erbium metal
doped As2S3 films using co-thermal evaporation. Using compounds like ErCl3 might be
a possible way to eliminate the clustering.
Table 5.2 Lifetimes of as-deposited Er3+:As2S3 films with ErCl3 as dopant.
Er3+ concentration in As2S3 film 1/e lifetime /
ms
Intrinsic
lifetime / ms
PL intensity /
mV
Low doped (0.2 at%) 2.5 2.6 ~10
High doped (0.4 at%) 1.4 1.6 ~2
Er3+ (metal):As2S3 (0.15 at% reference) 1.2 1.4 ~50
The PL intensity from these ErCl3 doped As2S3 films was relatively weak. The PL
intensity of the 0.2 at% ErCl3 doped sample was 1/5 and the 0.4 at% film was only 1/25
of that of the 0.15 at% erbium metal doped As2S3 film. In further investigation, a high
density of particulates with an average size of 10 μm diameter was found to cover the
whole film. An image of the particulates in ErCl3 doped As2S3 films taken in dark field
mode is shown in Figure 5.1. These particulates may act as defects and source of
absorption, and quench the PL emission immediately. It is also possible that these
particulates are ErCl3 condensate, meaning a proportion of the material never got into
the film leading to a low concentration of active Er3+ and thus a weak PL intensity.
More particulates were found on the 0.4 at% doped film than the 0.2 at% film,
suggesting that these particles are ErCl3 concentration related. To verify this
assumption, an As2S3 film without ErCl3 doping was deposited under the identical
conditions to the ErCl3 doped As2S3 films and was inspected again. The new As2S3 film
was almost particle-free under microscope. This indicated clearly that ErCl3 is the cause
of the particles. The mechanism of particle formation and the role of the ErCl3 in this
are still unknown, but the high density of particles make the films unusable for
waveguide amplifiers, and so no further investigations were performed.
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Figure 5.1 Particulates in co-thermal evaporated ErCl3 doped As2S3 films taken with an
optical microscope in dark field mode.
5.1.2 Co-thermal evaporation of Er2S3 with As2S3
Given the particulate issues described above, the alternative compound Er2S3 was
investigated. Unlike ErCl3, Er2S3 was widely used in chalcogenide glasses for Er3+
doping as it can easily be incorporated in the chalcogenide glass host matrix, especially
so in sulfur containing materials. For example, in Kasap et al.’s report [121], optical and
PL properties of powdered Er3+-doped GeGaS glasses of various mean sizes of near-
stoichiometric composition (Ge28Ga6.2S65.3:Er0.5 with 2.1x1020 ions/cm3 Er3+ ions), were
examined carefully, where Er2S3 was employed as the dopant source. The radiative
lifetime for the 4I13/2 energy level was calculated 2.6 ms using Judd–Ofelt theory, and
the estimated maximum emission cross-section at 1.53 μm was 15.5x10−21 cm2.
Tonchev et al. [295] reported both the thermal stability and optical properties of
(GeS2)75(Ga2S3)25 glasses doped with large amounts of Er2S3 (1.8 to 2.4 mol% of Er2S3).
Samples were synthesized with the melt-quenching method, and the measured ~ 1540
nm PL decay times were in the 1.13-1.55 ms range decreasing with increasing erbium
ion content. In Allen et al.’s work [119], besides ErCl3, in another set of GeGaS and
GeGaSs samples, Er2S3 was chosen as the dopant. Er2S3 doping concentration varied in
the range of 0.3-3 at%. PL spectra and decay lifetime of 1.54 μm emission were
measured under 980 nm excitation. The Er2S3-doped samples exhibited the 4I13/2 energy
level lifetimes in the 1-4 ms range. The PL efficiency measurements showed quenching
occurred at Er3+ concentration greater than 1 at%. Er2S3 was also used in erbium ion
doped film deposition. Lyubin et al. reported the co-thermal evaporation with Er2S3 to
fabricate Er3+-doped As2S3 and As2Se3 films [128]. The chalcogenide glasses were
evaporated from a quartz crucible, while a tungsten boat was used to evaporate the
Er2S3 powder on account of its high evaporation temperature. PL spectra at 1.54 μm of
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Er3+ doped As2S3 and As2Se3 films were recorded under Ar+-laser excitation, and a
linear PL intensity dependence on the exciting laser beam intensity was also observed.
In this research, the Er2S3 doped As2S3 film deposition procedure was similar to the
previous ErCl3 doped As2S3 film deposition procedure, which was described in Section
5.1.1, except replacing the ErCl3 from the resistively heated alumina crucible with high
purity Er2S3 (5N) powder. A test run of evaporating Er2S3 was commenced before the
Er2S3 doped As2S3 film deposition, to establish the Er2S3 evaporation rate-temperature
relationship. In the test run, even heating the alumina crucible up to its maximum
temperature of 1500 °C, there was still no sign of Er2S3 evaporation, indicating Er2S3
had a much lower vapour pressure than ErCl3. After cooling down the system and
venting the chamber, the alumina crucible loaded with Er2S3 was found to have been
damaged, as shown in Figure 5.2 (a). It is known that Al2O3 and phosphate are added in
silicate glasses to improve the dispersion of rare-earth ion in glass matrix. In this
experiment, the alumina crucible may react with Er2S3 which caused the damage of
alumina crucible [296-298].
Figure 5.2 Damaged Al crucible in Er2S3 evaporation (a); scanning electron image of an
as-deposited Er2S3 doped As2S3 film with tantalum boat (b).
Obviously, an evaporation boat for Er2S3 powder that can stand even higher
temperature (more than 1500 °C) was needed. Therefore, a tantalum boat was used in
the following deposition, and As2S3 was relocated to an alumina crucible. With this new
arrangement, evaporation of Er2S3 was observed from rate sensor when 80% of full
power (300 W) was applied to the tantalum boat.
However, in the following lifetime measurement, no 1.55 μm emission signal was
obtained from this film. The film surface was clearly rough under optical microscope
inspection. A scanning electron microscope (SEM) image of the film surface is shown
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in Figure 5.2 (b) with magnification of 2x104 confirming the poor surface quality. It is
speculated that the high temperature of the Er2S3 source (>1500 °C) radiatively heated
the wafer located 40 cm above the source. Prior unpublished experiments within the
Laser Physics Centre have established that at much above 70 °C wafer temperature, the
film quality of As2S3 films degrades rapidly. Three viable ways are proposed to resolve
this problem. Firstly, a cooling system could be employed to prevent the substrate wafer
temperature increase caused by heating of Er2S3 source. However, given that the wafers
are on a planetary style rotation carousel, this is a nontrivial implementation. Secondly,
the distance between the Er2S3 source and substrate wafer could be increased. In the
current chamber there is insufficient room to get more than approximately another 10
cm of throw distance, likely insufficient to resolve the issue given the ~10x3 cm size of
the source. Lastly, a water-cooled shield could be placed over the source leaving only
the emitting aperture exposed to reduce the heat load. Due to the complexity of the
deposition system, (6 thermal sources plus three sputter guns) there are no spare
feedthroughs available to be able to try this option. Consequently, no further progress
was made on thermally evaporated Er2S3 doped As2S3 film fabrication, and its viability,
as a method to fabricate high performance waveguide amplifier is still an open question.
5.2 Radio-frequency (RF) sputtering of erbium ion doped
As2S3 films
5.2.1 Introduction to RF sputtering chalcogenide glass films
RF sputtering of chalcogenide glasses using argon has been reported as a promising
method to deposit device-quality films [217, 279, 299, 300]. Typically, sputter
deposition is performed at pressures of 0.1-1 Pa, and involves creating a plasma (usually
in an inert gas like argon) by applying an RF voltage between a cathode, (which is the
target holder), and anode (which refers to the gun shield and the rest of the vacuum
chamber). With the bombardment by high energy ions, the target surface ejects atoms or
atom clusters that diffuse away and finally condense on the substrate wafer to form
films. The stoichiometry of the deposited films can be controlled through adjusting the
sputtering parameters, typically the chamber pressure, sputtering power and gas flow.
There are reports on rare-earth ion doped chalcogenide glass films fabricated by RF
sputtering. In Nazabal et al.’s report [300], Tm3+ doped Ga5Ge20Sb10S65 ( at% 2S2G)
films were fabricated using RF sputtering. The sputtering was performed in an Ar
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pressure of 3.7x10-2 to 3.7x10-3 Torr (~0.49-4.9 Pa), and the RF power varied between
20 and 50 W applied on a 2-3 mm thick 80 mm diameter target. The composition of the
films was found to be comparable to the bulk glass target with only a slight increase of
gallium. Lifetime of the 3H4 energy level of Tm3+ was measured under 800 nm
excitation and was found to be about 50 μs in a 0.1 at% doped sample, which was
shorter than the value of 120 μs from 2S2G powders with 0.05 at% Tm3+ doped
concentration. It was suspected that the presence of impurities or defects in the films
induced this discrepancy. Thin films of Er3+ doped As2S3 and Ge33As12Se55 (at%) have
also been fabricated using RF sputtering in Fuchs et al’s work [217]. Sputtering targets
were fabricated from Er3+ doped bulk samples (25 mm in diameter) and from
commercial un-doped targets (50 mm in diameter) with erbium metal pieces placed on
the target surface. The sputtering was performed at pressure of 3x10-2 Torr (~4 Pa), and
the applied RF power was 20 W. A PL lifetime of 4 ms of the 4I13/2 energy level was
measured on the obtained Er-doped As2S3 films under 977 nm excitation. However, in
the Ge33As12Se55 (at%) waveguides fabricated using wet etching, insertion losses at 1.3
μm (well out of the Er3+ absorption band) were larger than 30 dB, which was thought
due to the large porosity induced by column-like structure in the film which produced
strong light scattering.
5.2.2 Erbium ion doped As2S3 film deposited using RF sputtering
At the Laser Physics Centre of the Australian National University, the facility used for
co-thermal evaporation can also function as a sputtering chamber, containing three RF
sputtering guns. A 3-inch commercial As2S3 powder sintered target was used in the first
experiment, and a piece of erbium foil stuck onto the As2S3 target surface acted as an
erbium ion source. The chamber configuration can be found in Figure 3.5, and the As2S3
target with erbium foil attached on its surface is shown in Figure 5.3. Argon was used
for the creation of the sputtering plasma. During the deposition process, the substrate
was mounted on a 150-mm diameter rotatable holder. The glass target was loaded into
the 3-inch magnetron which was mounted 45° off vertical and facing upwards with a
distance of 100 mm between the target and substrate. Parameters such as Ar flow, RF
sputtering power and sputtering rate were optimised for film quality. The final
parameters used in the deposition are listed in Table 5.3.
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Figure 5.3 A 3-inch As2S3 RF sputtering target with a piece of erbium metal on the surface
acted as erbium ion source.
Table 5.3 Parameters used in erbium ion doped As2S3 film deposition using RF sputtering.
Base chamber pressure 1.5x10-7 Torr (~2x10-5 Pa)
RF power 43 W
Ar flow 20 sccm
Sputtering pressure 2 mTorr (~0.26 Pa)
Sputtering rate ~0.1 nm/s
Wafer rotation speed 20 r/min
5.2.3 Characterisation of RF sputtered erbium ion doped As2S3
The as-deposited RF sputtered Er3+:As2S3 films, coated with 100 nm SU-8 protective
layer, were characterised using the FilmTek. Film thickness, refractive index and
bandgap of the obtained films were measured at the film centre, as well as the film edge,
to check the uniformity of the whole film. This measurement was repeated on the same
film after 24 hours of 130 °C vacuum thermal post-treatment. The measurement results
of the as-deposited film, and the film after thermal treatment, are shown in Table 5.4.
Table 5.4 Physical properties of Er3+:As2S3 film deposited using RF sputtering.
Measured spot position on film / mm Thickness
/ nm
Bandgap
/ eV
RI @ 1550 nm
As-deposited, centre (0,0) 848 2.20 2.409
As-deposited, edge (35,0) 878 2.20 2.403
After thermal post-treatment, centre (0,0) 843 2.26 2.387
Compared with thermally evaporated films (see Section 3.2.3), RF sputtered films
have a lower bandgap (2.20 eV vs. 2.32 eV), but a higher refractive index close to the
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bulk value (2.41 at 1550 nm vs. of 2.35). The RF sputtered film thickness increases
from 848 nm (at film centre) to 878 nm (35 mm off centre), indicating that sputtering
led to less uniformity in films. This phenomenon could be explained partly by the
relatively short distance and the big angle between the target and the wafer, where the
edge of the wafer is closer to the source and so collects more condensate. The thermal
treatment applied has the opposite effect on RF sputtered films compared with what it
does on co-thermally evaporated films. In the RF sputtered films, after 24 hours thermal
post-treatment at 130 °C, the film bandgap increases up to 2.26 eV from 2.20 eV, while
the refractive index drops to 2.387 from 2.409 in as-deposited film, coupled with a
reduction in the thickness (the 6 nm reduction in thickness is certainly well within the
accuracy/resolution of the instrument and so is considered real). The differences
between RF sputtered and thermally evaporated films are thought to be micro-structure
related, thus careful attention should be paid to this when using such films in a real
device.
The lifetime and PL intensity were measured in as-deposited and thermal post-
treated (at 130 °C for 24 hours in vacuum) RF sputtered Er3+:As2S3 films, and the
results are shown in Table 5.5. A 1.3 ms 1/e lifetime of the 4I13/2 state of Er3+ was
achieved from the as-deposited film, which is longer than the 1.16 ms in the previous
0.15 at% co-thermally evaporated Er3+:As2S3 film in Section 3.6. After thermal post-
treatment, the 1/e lifetime of RF sputtered Er3+:As2S3 films increased to 1.9 ms, and the
PL intensity was enhanced roughly 4.5 times. Clearly, thermal treatment enables some
of the erbium ions to re-bond in a more ideal way reducing nonradiative transitions or
causes defect/impurity state densities to be significantly reduced.
Table 5.5 Lifetime & PL intensity of the 4I13/2 state of Er3+ in Er3+:As2S3 film deposited
using RF sputtering.
Er3+:As2S3 films of different
types
1/e lifetime / ms Intrinsic lifetime / ms PL intensity / mV
As-deposited 1.3 1.6 40
After annealing 1.9 2.1 180
Ref: Thermal deposition (0.15
at%) (see Section 3.6)
1.16 1.35 ~300
As mentioned in Section 5.2.2, to dope erbium ion into the As2S3 films, a piece of
erbium foil was stuck on the As2S3 target surface. Due to the different sputtering rates of
the erbium foil and the As2S3 target, it was not possible to estimate the erbium ion
concentration in the final Er3+:As2S3 films based on the proportion of areas each
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material exposed to plasma. The available energy dispersive X-ray spectroscopy (EDX)
also could not resolve clearly the erbium ion concentration at the 0.5 at% level and so
was not of assistance.
Another potential method to estimate the Er3+ concentration in Er3+:As2S3 film, is to
compare its Er3+ absorption band intensity with a film of known Er3+ concentration in
the same host. In the absence of clustering, the Er3+ absorption is proportional to its
concentration [198, 301]. Prism coupling offers a widely used way to measure the
absorption intensity of a film. However, good film quality is critical to getting an
accurate absorption or propagation loss in prism coupling methods described in Section
3.3.2(c). Inspection under a microscope showed micro-sized particles with diameter
ranging from 1 μm up to 4 μm covering the whole wafer at a much higher density than
usual thermally evaporated As2S3 films. A micrograph of the particles on the RF
sputtered Er3+:As2S3 film taken in dark field mode with an optical microscope is shown
in Figure 5.4. Given these particles would cause light scattering and an indeterminate
propagation loss contribution, this measurement was not pursued.
Figure 5.4 Particulates on RF sputtered Er3+:As2S3 film, which was taken with microscopy
in dark field mode. The micro-sized particulates have diameters ranging from 1 μm up to
4 μm.
It was thought, based on prior experience with other materials, that the particulate
problem originated from the sintered powder nature of the As2S3 target and the visible
cracks and graininess covering the whole target surface after use (see Figure 5. 3).
Therefore, another 3-inch one-side polished As2S3 glass disc was purchased as a new
target. Er3+:As2S3 films with better quality (checked using microscope in dark field
mode) were deposited under same conditions as previously described in Section 5.2.2
with this new As2S3 glass target. The lifetime and PL intensity of the as-deposited and
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thermally treated Er3+:As2S3 films were measured, and the results are shown in Table
5.6.
Table 5.6 Lifetime and PL intensity of Er3+:As2S3 film deposited using RF sputtering with
the new As2S3 glass target.
Er3+:As2S3 films of different
types
1/e lifetime / ms Intrinsic lifetime / ms PL intensity / mV
As-deposited 1.18 1.45 ~120
After annealing (130 °C for
24 hours in vacuum)
1.82 1.97 ~440
Ref: Thermal deposition
(0.15 at%) (see Section 3.6)
1.16 1.35 ~300
The lifetime of the as deposited film using As2S3 glass target is at the same level as
the previous films sputtered with the sintered As2S3 target, but the PL intensity
increases significantly: from 40 mV to 120 mV for the as-deposited film, and 180 mV to
440 mV after thermal post-treatment, which is even stronger than the best co-thermal
evaporated Er3+:As2S3 film with 0.15 at% erbium ion described in Section 3.6. The
erbium ion concentration in this waveguide is uncertain, but as will be discussed below
it is subsequently estimated at 0.1 at%.
5.2.4 Pump-induced PL intensity decay in RF sputtered erbium ion doped As2S3
films
It was noticed that the PL intensity of RF sputtered Er3+:As2S3 films dropped
significantly during PL testing, and fell to almost zero after 48 hours of green light
treatment (light wavelength of 550 nm with intensity of ~10 mW/cm2 from a typical
LED flood light in ambient air), the normal light treatment for photoinduced annealing
of our As2S3 films. Interestingly, the reduced PL intensity of this RF sputtered
Er3+:As2S3 film could be restored to its original value if this sample was then thermally
treated (130 °C in vacuum for 24 hours). It was also noticed that this process was
repeatable, and the PL intensity cycled with different post-treatment types as shown in
Figure 5.5. The T in x-axis refers to 24 hours thermal post-treatment under vacuum, and
the L refers to 48 hours green light post-treatment (light wavelength of 550 nm with
intensity of ~10 mW/cm2) in ambient air.
123
Figure 5.5 PL intensity cycled with light and thermal post-treatments of RF sputtered
Er3+:As2S3 film. T refers to 24 hours thermal post-treatment under vacuum and the L
refers to 48 hours green light post-treatment (light wavelength of 550 nm with intensity of
~10 mW/cm2) in ambient air.
To characterise the PL intensity change of RF sputtered Er3+:As2S3 films with time
in thermal and light post-treatment, the PL intensities of two Er3+:As2S3 film samples
(#1 & #2) were measured over time during sequential light & then thermal treatment
processes, and the results are shown in Figure 5.6. During light treatment, the PL
intensity of the Er3+:As2S3 film dropped significantly in the first few hours and then
slowed down to reach almost zero in about 60 hours. In subsequent thermal treatment,
the PL intensity of the light treated Er3+:As2S3 film increased quickly and reached its
maximum value in ~20 hours.
Figure 5.6 PL intensity of a RF sputtered Er3+:As2S3 film drops in time during light
post-treatment (a), and the PL intensity of the light post-treated Er3+:As2S3 film increases
in time during thermal post-treatment (b).
It was suspected that the green light and moisture in the atmosphere were
responsible for the PL intensity change in Er3+:As2S3 films. To figure out exactly what
factors caused this phenomenon, 5 pieces of RF sputtered Er3+:As2S3 film cleaved from
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the same wafer were put in 5 different conditions, and the PL intensity was monitored
with time. The five different conditions were as follows:
1) Sample immersed in deionised water (DI water);
2) Sample immersed in DI water and irradiated with green light;
3) Set sample in ambient air and irradiated with green light;
4) Sample irradiated with green light in an N2 atmosphere;
5) Set sample in a high vacuum chamber (~1x10-7 Torr, equals to ~1.33x10-5 Pa) and
irradiated with green light through a glass window.
The green light in this experiment was the same light used in previous light post-
treatment (centre wavelength at 550 nm with intensity of ~10 mW/cm2).
From the results shown in Figure 5.7, Er3+:As2S3 films in condition No.1 (DI water
only) and No.5 (under green light irradiation in high vacuum) keep the PL intensity
without obvious decay for a long time. PL intensities of Er3+:As2S3 films in other
conditions dropped significantly and the decay in sample No.2 (immersed in DI water
with green light irradiation) is the most significant one. Therefore, it is clear that water
(or moisture in air) and green light (550 nm) are essential for the PL intensity decay
described in the beginning of this section. One may argue that the sample No.4 (N2 flow
with green light irradiation) is not supposed to decay that fast, because the moisture in
air should be expelled by N2 flow. However, in the test setup, chamber used to keep the
sample and inject the N2 flow in was not sealed perfectly, thus it is believed there was
still moisture in that chamber which was sufficient to cause this PL intensity decay.
Figure 5.7 PL intensity of RF sputtered Er3+:As2S3 films change in time under various
storage conditions. DI water is deionised water.
125
To further enhance the evidence of water diffusion in waveguide, a undoped RF
sputtered As2S3 film (no erbium ion contained in this film) was patterned into
waveguides, and the propagation loss spectrum of a 2 μm wide waveguide was recorded
after each thermal (noted as T in Figure 5.8) and green light post-treatment (noted as L
in Figure 5.8), repeatedly, and is shown in Figure 5.8. An evident absorption band at
1430-1480 nm shows up after each light post-treatment and disappears after each
thermal post-treatment. An absorption band at this wavelength is considered to be
associated with the overtones of the O-H stretch vibration near 3600 cm-1 [302], which
supports the assumption that moisture in air is one of the main factors contributing to
this phenomenon.
Figure 5.8 Loss spectra of a RF sputtered As2S3 waveguide after thermal/light post-
treatment. These spectra were recorded with OSA and using SC source as probe source.
In this figure, T refers to the thermal post-treatment and L refers to the light post-
treatment, and the ‘T→L’ refers to a thermal post-treatment followed by another light
post-treatment.
In order to get the erbium ion absorption spectrum in RF sputtered Er3+:As2S3 film,
a RF sputtered Er3+:As2S3 film was patterned into waveguides, and the recorded
spectrum of a 3 μm wide waveguide (the same structure as depicted in Figure 4.7 except
the width changed from 2 μm to 3 μm) without post-treatment is shown in Figure 5.9.
Absorption bands related to O-H stretch vibration and erbium ion are clear and overlap
to each other at shoulder. From Figure 5.8, the O-H stretch vibration absorption band
ends by 1520 nm, therefore it is still possible to get the erbium ion absorption value at
wavelengths longer than that. The peak erbium ion absorption at 1540 nm was 6 dB for
a 48 mm long waveguide, which results in absorption of 1.2 dB/cm. Compared with the
previous number of 1.9 dB/cm from the 0.15 at% Er3+:As2S3 waveguide deposited using
co-thermal evaporation (see Section 4.5.1), the erbium ion concentration in this film is
estimated about 0.1 at%.
126
Figure 5.9 Spectrum of erbium ion absorption overlaps with O-H stretch vibration
absorption of a 48 mm RF sputtered Er3+:As2S3 waveguide.
There has been observation of degradation of PL intensity in Er3+ doped
chalcogenide film deposited using RF sputtering. In Fuchs et al.’s reports [217], Er3+
doped As2S3 and Ge33As12Se55 (at%) chalcogenide films were fabricated using RF
sputtering. Column-like structure was found in Er3+:Ge33As12Se55 waveguide. No PL
signal was observed from it, which was believed to due to the huge loss of the
waveguide. It was also noticed that the PL intensity in the as-deposited Er3+:As2S3 film
was weak, but after thermal post-treatment, the PL intensity increased up to 40 times.
Large PL intensity increase after thermal post-treatment is similar to observations in this
research. Fuchs explained the absence of PL in Er3+:Ge33As12Se55 waveguides was due
to light diffusion from columnar structure in the films. However, in the current
experiment, columnar structure based light scattering cannot explain the PL intensity
change in the RF sputtered Er3+:As2S3 film after thermal and light treatment; also it
cannot explain the appearance of the O-H stretch vibration absorption band in the RF
sputtered As2S3 waveguide after green light treatment. Furthermore, it is clear that any
columnar growth in the film must be densely packed as the waveguide losses are low,
where the propagation loss is estimated ~0.4 dB/cm in this waveguide.
To explain the current measurement results, a modification was made on Fuchs’s
explanation. In green light post-treatment, due to the existence of nano columnar
structure in RF sputtered Er3+:As2S3 film, moisture from the air could easily diffuse into
the film along the column boundaries or other types of nano pores resulting from the
film growth. As the energy of the 550 nm green light (2.25 eV) is above the bandgap of
Er3+:As2S3 film (~2.2 eV), it has the capability to induce bond rearrangement and
therefore potentially to form O-H bonds with film constituents. It is well know that O-H
127
quenches erbium ion emission effectively [152], thus the PL intensity in Er3+:As2S3 film
drops, also the appearance of O-H absorption band after light post-treatment could be
explained by this hypothesis. In thermal post-treatment, bonds between O-H and
Er3+:As2S3 are broken down with the help of thermal energy, and moisture is expelled
from the film under vacuum, and thus the PL intensity recovers back to its high level
and the O-H absorption band disappears in the RF sputtered Er3+:As2S3 waveguide.
An obvious question raised from here is whether the decay of PL intensity observed
under green light in air affects Er3+:As2S3 waveguide amplifier performance. Excitation
experiments on RF sputtered Er3+:As2S3 waveguides were carried on in ambient air.
Although no green light was irradiated directly on the waveguide in excitation
experiments, green light due to up-conversion was observed in the Er3+:As2S3
waveguide.
From the measured wideband PL spectrum from the erbium ion doped As2S3 co-
thermal waveguides (see Figure 4.19), there are three up-conversion related bands that
could act as ‘green light source’ here. The first one is centred at 530 nm (2H11/2→4I15/2),
the power level of this emission band is at 4x10-14 W with ~20 nm bandwidth based on
the OSA recorded spectrum, and this data was taken with a 1 nm resolution bandwidth;
the second one is at 550 nm (4S3/2→4I15/2) with 6x10-14 W power level, ~30 nm
bandwidth; the last one is 660 nm emission (4F9/2→4I15/2) with 1.1x10-13 W power level,
~40 nm bandwidth. Summing all energy together and considering the effective mode
area of the well confined waveguide is ~1 μm2, the real ‘green light’ intensity is
estimated ~1 mW/cm2, almost the same order of intensity as the green light used for
light post-treatment. Also, all data used here is from the recorded spectra, therefore the
real green light intensity in the waveguide should be ~3 dB (0.9 dB for one side fibre-
to-fibre from the current optical set-up, and 2.35 dB for fibre to waveguide mode
mismatch and surface reflection) stronger than the above value.
To verify this assertion, the ASE spectra of a 48 mm long RF sputtered Er3+:As2S3
waveguide were measured under 1490 nm excitation three times: immediately after the
pump laser was turned on, and 5 and 10 minutes after the pump laser was turned on.
The pumping intensity was ~2000 kW/cm2, and the results are shown in Figure 5.10.
About 3 dB drop of the ASE level is observed in the first 5 minutes, and another 1.5 dB
drop in the following 5 minutes. This decrease of ASE level indicates similar effects as
those seen in the PL measurements are being induced by upconverted green light, and
this will seriously affect amplifiers & laser devices.
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Figure 5.10 ASE spectrum of a 48 mm long RF sputtered Er3+:As2S3 waveguide excited
at 1490 nm.
Columnar structure is found in films fabricated using RF sputtering in many
materials [303, 304], as well as in chalcogenide glass films [217, 305]. Coating a
moisture barrier layer on the film to protect it is one option to resolve this issue, but
reducing the columnar structure in the film would be a better solution. Either or both are
essential for future research on erbium ion doped RF sputtered films, but neither was
available at the time.
5.3 Conclusion
In this chapter, different approaches to realising erbium ion doped As2S3 waveguide
amplifiers with reduced clustering were investigated. Erbium ion containing
components (ErCl3 and Er2S3) were chosen to replace the erbium metal in co-thermal
evaporation. With ErCl3, a long 1/e lifetime of the 4I13/2 state of 2.6 ms was achieved in
a low Er3+ doped (0.2 at%) sample, but the high density of 10 µm sized particles made
the films unusable for waveguide amplifiers. The evaporation of Er2S3 was unsuccessful
due to its high evaporation temperature which was beyond the upper limit of the
available deposition system using the shielded alumina crucibles and with a tantalum
boat resulted in very poor film quality. In RF sputtered Er3+:As2S3 film, a 1/e lifetime of
the 4I13/2 state up to 2.1 ms was achieved in a thermally treated sample, but film quality
became an issue that prevented the fabrication of suitable quality waveguides. With a
glass As2S3 target, film quality was improved and the PL intensity from the obtained
Er3+:As2S3 film (after thermal post-treatment) was even stronger than that of the film on
which internal gain was achieved in Section 4.5.3. However, due to the columnar
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growth of RF sputtering films water penetration and bonding to the glass matrix with
the help of above band gap light quenched the PL rapidly. Whilst none of the above are
fundamental issues and the results are promising, it was not possible to resolve these in
the time available and with the equipment available.
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131
Chapter 6
Erbium ion doped Ge-Ga-Se waveguides
Although internal gain was achieved in an erbium ion doped As2S3 waveguide amplifier
as reported in Chapter 4, net gain in a rare-earth doped chalcogenide waveguide
amplifier is still out of reach because of the clustering and the inability to disperse the
clusters through thermal post-treatment due to the low intrinsic rare-earth ion solubility
of As2S3. In this chapter, a new chalcogenide glass host, the Ge-Ga-Se system, which is
known to have high rare-earth ion solubility, is investigated. Bulk glasses, thin films
and waveguides based on erbium ion doped Ge-Ga-Se material were fabricated and
characterised. Record erbium ion population inversion was achieved in these
waveguides. But excitation laser upconversion pumped photoinduced absorption was
encountered which presented a new hurdle to realising net gain.
6.1 Erbium ion doped Ge-Ga-Se bulk glasses
6.1.1 Background on Ge-Ga-Se glasses
Internal gain between 1570-1630 nm was achieved in an erbium ion doped As2S3
waveguide amplifier [133] in the work detailed in Chapter 4. The total erbium ion
population inversion was, however, estimated only at ~30% from the absorption and
enhancement curves, and this was still some way from realising a useful amplifier with
net gain. Analysis revealed that thermally evaporated erbium ion doped As2S3 films
showed clear signs of erbium ion clustering even at concentrations as low as 0.15 at%
(~4-5×1019 ion/cm3) (see Section 3.6 and 4.5.3). This erbium ion concentration is
essentially the minimum usable doping level for a planar waveguide device. Under
1490nm pumping, the maximum inversion is 65% which limits the maximum possible
internal gain to ~0.5 dB/cm, even before waveguide losses, ion-ion effects, etc. are
considered which further reduce the gain. Long waveguides with lower doping are not a
solution here, as the passive losses at the pump wavelength then become a problem
(passive losses of ~3-5 dB are tolerable for the pump which limits amplifier length to
perhaps 15 cm for the ~0.3 dB/cm propagation losses common for high refractive index
contrast waveguides of this type [62]). Obviously, significant improvement is necessary
to realise a useful high gain amplifier. Different approaches to doping erbium ions into
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As2S3 films for better erbium ion homogeneity without thermal post-treatment, or the
use of hosts with better intrinsic erbium ion solubility are considered promising
solutions to this issue. The former was presented in Chapter 5. This chapter focuses on
alternative hosts.
Among the family of chalcogenide glasses, gallium containing materials are known
to have good rare-earth ion solubility. In As2S3, the three dimensional network can be
interpreted using a [AsS3/2] trigonal pyramidal units. The insufficient quantity of non-
bridging sulfur in As2S3 that is needed to coordinate the isolated rare-earth ions, cannot
meet the requirement of the large coordination number rare-earth ion needs to fit in,
therefore, resulting in low rare-earth ion solubility [140, 306]. The incorporation of
gallium as modifiers into chalcogenide glassy networks could dramatically increase the
solubility of rare-earth ions due to the presence of edge-sharing [GaS4] (or [GaSe4] in
selenium glass) tetrahedral structure, which provides a compensation for the negative
charge of free S2− (or Se2- in selenium glass) ions by forming the chemical bonds with
rare-earth ions [307-309].
Erbium ion doped Ga-containing chalcogenide bulk glasses were intensively
investigated, and promising 1540 nm emission properties were obtained, e.g. [119, 198,
199]. Ikuta et al. measured PL intensity at 1540 nm in Er3+ doped stoichiometric
(GeSe2)1–x(Ga2Se3)x chalcogenide glasses and concluded that Er3+ could be well
dispersed and optically active up to 2 at% [198]. The importance of gallium for Er3+
activation was also studied, and the PL results indicated the critical gallium to erbium
ion concentration ratio was about 5. Too much gallium in the glass may cause
inhomogeneity. Allen et al. studied the optical characteristics of a series of Er3+-doped
chalcogenide glasses and found that under 980 nm excitation, the shortest 1550 nm
emission decay time of 1–1.5 ms were observed in GaGeAsSe glasses, whilst the
highest values of 2–4 ms were obtained in GaGeSe and GaGeS samples [119]. Koughia
et al. examined the optical properties of erbium ion doped Ge-Ga-Se and Ge-Ga-S
chalcogenide glasses [199]. The lifetime of the 4I13/2 state of Er3+ in Ge-Ga-Se glass was
predicted at 1.8 ms from Judd-Ofelt theory, but in experiments, with different
excitation wavelengths (976 nm, 818 nm and 662 nm), longer lifetimes ranging from 2.7
to 3.1 ms were observed in a 2 at% erbium ion doped sample, and this discrepancy was
explained by effective energy transfer in the Er3+ system when the self-absorption effect
and the possible re-emission and re-absorption processes were taken into consideration.
Also in another report, in Ge-Ga-Se glass system, it was suggested that a high ratio of
133
GaIII to Er3+ ensured homogeneous distribution of Er3+, whilst a low ratio led to the
formation of clusters [120]. In the intermediate region, energy was shown to be able to
efficiently migrate from one Er3+ to another. Tonchev et al. doped Er2S3 from 1.8 to 2.4
mol% into Ge-Ga-S glasses [295], and a broad PL emission band at ~1540 nm with
decay time in the range of 1.13 to 1.55 ms was observed under 818 nm excitation. The
PL decay time decreased rapidly with increasing Er3+ concentration in the range of
study, at a linear rate of 0.7 ms/mol%. All of these works showed the potential of Ge-
Ga-Se glass as a promising host for Er3+ doping due to the long 1.5 μm PL decay time
and high dopant concentration possible.
To gain an initial understanding of the properties of the Er3+:Ge-Ga-Se glass system,
six pieces of Er3+:Ge25Ga10Se65 (at%) glass with different Er3+ concentrations were
prepared. Absorption and emission properties of the obtained Er3+:Ge25Ga10Se65 glasses
were investigated in order to determine the effect of Er3+ concentration and establish a
baseline for the performance of Er3+ in these hosts to compare with thin films.
6.1.2 Erbium ion doped Ge-Ga-Se glasses synthesis and characterisation
In order to provide a benchmark for film based materials, erbium ion doped
Ge25Ga10Se65 (at%) chalcogenide glasses with different Er3+ concentrations (ranging
from 0.1 at% (~3×1019 ion/cm3) to 2 at% (~6×1020 ion/cm3)) were prepared from high
purity (5N) germanium, gallium, selenium and erbium metals using the conventional
melt-quenching method. This composition (Ge25Ga10Se65 (at%)) was chosen as it was
expected to be a stoichiometry that should produce thin films with properties close to
the bulk glass [225] and that had sufficient gallium to enable ~1 at% erbium ion to
remain unclustered. The required amounts of these raw materials were weighed inside a
dry nitrogen glove box and loaded into a pre-cleaned quartz ampoule. The ampoule was
then sealed under vacuum (~10-4 Pa) using an oxygen-hydrogen torch, and introduced
into a rocking furnace for melting of the contents at 950 °C. The melt was homogenized
for a period not less than 12 hours, then the ampoule was removed from the rocking
furnace and water quenched. The resulting glass boule was subsequently thermally
treated at a temperature 20 °C below its glass transition temperature Tg of ~395 °C for 2
hours, before being slowly cooled down to room temperature. The obtained glass rods
were cut into small pieces with ~2 mm thickness and then polished on both surfaces.
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The room temperature Raman spectra of these Er3+:Ge25Ga10Se65 (at%) bulk glass
samples were measured using a Horiba Jobin Yvon 64000 spectrometer employing a
632 nm laser as the excitation source. The 830 nm excitation source was avoided,
because this wavelength could excite the erbium ions and thus lead to emissions which
may overlap with the Raman signal. A CCD array detector was installed on the
spectrometer to record Raman spectra. Raman spectra collected from erbium ion doped
Ge25Ga10Se65 (at%) bulk glass samples are shown in Figure 6.1 with vertical offsets to
aid viewing.
Figure 6.1 Raman spectra of Er3+:Ge25Ga10Se65 (at%) bulk glasses with Er3+
concentrations ranging from 0.1 at% to 2 at%, under excitation at 632 nm.
The main feature in the Raman spectra is the band with its maximum near 202 cm-1
with two shoulders near 177 and 217 cm-1. According to reference [310], the strongest
band rising near 202 cm-1 is attributed to the v1(A1) symmetric stretching modes of
corner-sharing [GeSe4/2] tetrahedra. The shoulder at 217 cm-1 of the dominant band can
be associated with the Ac1 breathing vibrations of distorted fragments of the layered c-
GeSe2 structure containing edge-shared Ge2Se8/2 bi-tetrahedra. The Raman band at 177
cm-1 is probably related to stretching modes of Ge-Ge, Ga-Ga or Ge-Ga bonds (AG
vibration modes) in Ge2Se6/2, Ga2Se6/2 or GeGaSe6/2 structural units. This 177 cm-1 band
is much stronger in the 2 at% doped sample implying a significantly increased
concentration of Ge-Ge, Ga-Ga homopolar bonds in this glass [311]. Although the band
at 177 cm-1 in the 2 at% sample is evident, no sign of the evolution of this band in the 1
at% and below samples is visible, and the cause for this abrupt change is not clear.
Besides the obvious Raman bands discussed above, the broad Raman band of lower
intensity (235-330 cm-1), consisting of at least three weak bands with strong overlaps,
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can be assigned to the vibrations of Se-Se bonds (245, 265 cm-1) and to the vibrations of
GeSe4 structural units (300 cm-1). Additionally, for Ge-Ga-Se based chalcogenide
glasses, the vibration modes of Ga-containing structural units, such as [GaSe4] and
[Se3Ga-GaSe3] are usually overlapping with Ge-containing structural units [GeSe4] and
[Se3Ge-GeSe3], due to the small difference between bond strength and atomic weights
of Ga and Ge [310]. From Figure 6.1, the increasing of the erbium ion concentration
from 0.1 at% up to 1 at% brought almost no significant change in Raman spectrum and
therefore glass structure, whilst in the 2 at% Er3+ doped sample, a sign of more Ge-Ge,
Ga-Ga homopolar bonds appears, indicating a local structure change occurs in this high
doped sample. This is consistent with reports that the introduction of transition metals
into the glass matrix may cause significant change of almost all glass parameters, but
with rare-earth ions only mild change could be found [312, 313].
6.1.3 Optical properties of erbium ion doped Ge-Ga-Se glasses
Absorption spectra of Er3+:Ge25Ga10Se65 (at%) bulk glasses were recorded using a Cary
5000 UV-Vis-NIR spectrophotometer. The absorption curve of a 2.9 mm optical path
length, 0.5 at% Er3+:Ge25Ga10Se65 sample, is shown in Figure 6.2 for example.
Figure 6.2 Absorption spectrum of a 2.9 mm optical path length 0.5 at% Er3+:Ge25Ga10Se65
bulk glass.
Three absorption bands show up in the spectrum in Figure 6.2 as expected. The
strong band located at 1538 nm is due to the 4I15/2→4I13/2 transition. The 980 nm feature
corresponds to the absorption between the 4I15/2 and the 4I11/2 state, and the band at 805
nm is assigned to the 4I15/2→4I9/2 transition. The increasing absorption starting from
~800 nm to short wavelengths is the weak absorption tail (WAT) that is formed by
extrinsic impurities and intrinsic defects [39]. At even shorter wavelength the absorption
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increased exponentially as the band gap is reached, these effects preventing us
observing the ground state absorptions of higher energy levels.
Absorption features at 1538 nm from all 6 samples with different erbium ion
concentrations were extracted and are shown in Figure 6.3(a). The absorption spectra
shape are essentially invariant except the intensity increases with increase of erbium ion
concentration. The peak absorption value at 1538 nm was extracted and is plotted versus
erbium ion concentration in Figure 6.3(b), from which a linear relationship (fitted line)
between absorption and doped concentration up to 2 at% (~6×1020 ion/cm3) is seen
clearly, implying the erbium is not forming physical clusters at concentrations up to 2
at%.
Figure 6.3 Absorption band centred at 1538 nm for different erbium ion concentrations in
Er3+:Ge25Ga10Se65 bulk glasses (a); peak absorption value at 1538 nm versus erbium ion
concentration in Er3+:Ge25Ga10Se65 bulk glasses (b).
6.1.4 PL and PL lifetime of erbium ion doped Ge-Ga-Se bulk glasses
The room temperature photoluminescence spectra (PL) of the Er3+:Ge25Ga10Se65 (at%)
bulk glass samples were measured using the Horiba Jobin Yvon 64000 spectrometer
employing an 830 nm laser as the excitation source. An InGaAs detector installed on the
spectrometer recorded PL spectra up to 1600 nm. The 4I13/2 excited state lifetime of Er3+
in Er3+:Ge25Ga10Se65 (at%) glass was measured under 1490 nm excitation, using the all
fibre confocal arrangement previously described in Section 3.3.2. With 1490 nm
excitation, the effect of the possible transitions involved in higher excited states on the
measured 4I13/2 level lifetime was reduced significantly. PL intensity values were also
collected in the lifetime measurement set-up using both 1490 and 980 nm pumps.
Figure 6.4 shows the PL spectra of the bulk glasses with different erbium ion
concentrations excited with an 830 nm continuous wave laser. All the spectra have
similar line shape and band positions. The emission peak lays at 1538 nm, which is
137
slightly red-shifted from the 1532 nm usually observed in oxide glass hosts [314], also
was observed in Er3+:As2S3 films (see Section 3.3.2(c)). The long wavelength shoulder
on the PL spectrum is growing slightly with the increase of Er3+ concentration, which is
also plotted in the inset for the normalized PL spectra. Shoulder growth accompanied by
increasing emission bandwidth has previously been observed in a number of studies,
e.g.[120], and is believed to be associated with reabsorption of the emitted 1540 nm
light in highly doped materials. In this work, however, there is no sign of bandwidth
increase, which suggests a different mechanism is in play, perhaps relates to the higher
order ~1550 nm transition (4H11/2→4I9/2 transition), possible due to the 830 nm pump
wavelength.
Figure 6.4 PL spectra of Er3+:Ge25Ga10Se65 (at%) bulk glasses with different Er3+
concentrations ranging from 0.1 at% to 2 at% excited at 830 nm. The insert is the
normalised PL spectra.
The PL intensity at 1538 nm versus erbium ion concentration for the 830 nm pump
is presented in Figure 6.5(a). It shows an initial supra-linear increased in PL with
erbium ion concentration before becoming sub-linear at high concentrations, as would
be expected from ion-ion effects. Figure 6.5(b) shows the PL intensity with a 1490 nm
pump laser which displays the normally expected behaviour of reducing PL efficiency
with increasing Er3+ concentration, as seen with a 980 nm pump. The linear relationship
between absorption and erbium ion concentration (see Figure 6.3(b)) verifies that the
supra-linear curvature in the 830 nm PL result is not due to errors in the Er3+
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concentration at low values. The origin of this effect remains uncertain and needs
further investigation. However, it is clear that the performance is beginning to decline at
Er3+ concentration beyond 1 at%, which for this glass composition could be considered
optimal on the basis of the data presented. While considering the 10 at% Ga
concentration in all the glasses and the optimal ratio of GaIII/Er3+ at 10:1 noted in prior
work [120], the limit of the homogeneous distribution of erbium ion in this host should
be ~1 at%. Therefore it is reasonably expected that erbium ions start to form clusters at
an erbium ion concentration more than 1 at%, leading to the gradual quenching of PL
emission at 1550 nm [315, 316] which fits with the measured data.
Figure 6.5 PL intensity at 1538 nm with 830 nm excitation (a), and PL intensity at 1538
nm excited at 1490 nm (b) of Er3+:Ge25Ga10Se65 (at%) bulk glasses with erbium ion
concentration ranging from 0.1 at% to 2 at%.
6.1.5 Absorption and emission cross-section of erbium ion doped Ge-Ga-Se bulk
glasses
Based on the absorption spectrum of the 0.1 at% (~3×1019 ion/cm3) Er3+:Ge25Ga10Se65
(at%) glass sample, McCumber theory [161] was employed to calculate the emission
cross-section, and the results are shown in Figure 6.6. The McCumber theory and
measured emission cross-section agree well except for the minor deviation in the 1520
nm tail region. Also in Figure 6.6, a maximum pump efficiency curve is plotted which
represents the maximum inversion possible versus pump wavelength as described by
equation 2.14. With 1490 nm excitation, the inversion efficiency is around 71%. At
shorter wavelengths, the efficiency is higher, but the absorption cross-section reduces
further leading to lower overall pump absorption. The results also indicate that for a
highly doped short waveguide amplifier, the optimum pump wavelength lays in the
1500-1510 nm region where the absorption is approaching half its peak value whilst the
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inversion efficiency remains above 60%, as has been successfully employed in for
example TeO2 based waveguide amplifiers [152]. This is a significantly longer pump
wavelength than in silica or other oxide based hosts.
Figure 6.6 Normalised absorption, emission spectra and simulated emission cross-section
based on M-C theory, and max pump efficiency with pump at 1490 nm of the 0.1 at%
Er3+:Ge25Ga10Se65 bulk glass.
The dependence of PL intensity on pump power was investigated for high and low
erbium ion doped samples (0.1 at% (~3×1019 ion/cm3) and 2 at% (~6×1020 ion/cm3)),
with excitation at 1490 nm, and the results are shown in Figure 6.7. It is clear in both
samples that PL intensity is increasing with the increase of pump intensity, but with
different trends. For the 0.1 at% Er3+ doped glass, the PL intensity increases linearly
with pump intensity in the measured range which is low enough in intensity that there is
still a significant fraction of Er3+ in the ground state and no saturation of the emitted
1550 nm radiation occurred. For the 2 at% Er3+ doped sample, there is instead an
approximately quadratic relationship over the same pump power range. This quadratic
trend in the highly doped glass results from radiative and non-radiative ion-ion energy
exchange interactions which are not present in the lower doped samples due to the
larger average distance between Er3+ ions.
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Figure 6.7 PL intensity at 1538 nm versus pump intensity (at 1490 nm) in low (0.1 at%)
and high (2 at%) Er3+ doped Ge25Ga10Se65 bulk glass.
Due to the low phonon energy in chalcogenide glasses, the multiphonon relaxation
rate of Er3+ between the 4I11/2 state and the 4I13/2 state is low and the 4I11/2 state has a long
lifetime, comparable to that of the 4I13/2 state [46]. Therefore, the lifetime of the 4I13/2
state when pumped with a 980 nm laser would be expected to be significantly longer
than that with a 1490 nm laser, due to energy storage at the 4I11/2 state. To avoid this
issue, a 1490 nm laser was employed to probe the 4I13/2 state lifetime of the Er3+, using
the confocal experimental configuration described in Section 3.3.2, to avoid
reabsorption and stimulated emission artifacts [152]. It is well known that as the Er3+
concentration increases that the average distance between neighbouring ions diminishes,
enabling a variety of energy transfer effects to become relevant. As many of these
effects occur between two or more Er3+ ions in the 4I13/2 metastable state, these excited
ions do not contribute to creating photons in the desired state. Therefore, this reduces
the PL efficiency and also results in a decrease of the lifetime of the desired state with
increasing ion concentration, as the energy transfer is a fast ion-ion interaction.
Typically these effects have been found to follow an empirical formula first proposed in
[189]:
𝜏𝑜𝑏𝑠 =𝜏0
1+(𝜌 𝑄⁄ )𝑝 (6.1)
where τobs is the observed fluorescence lifetime, τ0 is the ideal fluorescence lifetime with
zero concentration, ρ is the Er3+ concentration, Q is the quenching concentration and p
referred as phenomenological parameter characterizing the steepness of the
corresponding quenching curve. As several of the effects encapsulated in this
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relationship also depend (sometimes) nonlinearly on the number of excited ions, there is
also a dependence of the parameters on pump power.
The photoluminescence decay curves were measured for the glass samples over a
range of different pump intensities (from ~10 to ~600 kW/cm2), the parameters recorded
being the measured 1/e lifetime, and the intrinsic lifetime (ie the supposed radiative
lifetime, see Section 3.3.2). The lifetimes versus pump power for each glass were then
fitted with suitable polynomials to extrapolate the data back to ‘zero pump power’ to
enable the effective ‘zero pump power lifetimes’ to be extracted. Figure 6.8 presents the
data for both the 1/e and intrinsic lifetimes of the 4I13/2 metastable state of Er3+ at
effective ‘zero power’ and at ~600 kW/cm2 pump intensity.
Figure 6.8 1/e and intrinsic lifetime of the 4I13/2 metastable state of Er3+ with different
doping concentrations, at extrapolated zero pump power (a); also at high pump power
(~600 kW/cm2 intensity) (b), of Er3+ doped Ge25Ga10Se65 bulk glass.
There are a number of clear trends from the data in Figure 6.8. Looking first at the
intrinsic lifetime in Figure 6.8(a) & (b), it is independent of the pump power (as
expected) and had a clear linear dependence on concentration with a slope of -0.48
ms/at% erbium ion. The slope of the line is also significantly lower than the -0.7
ms/mol% (Er2S3) reported in [295]. However this improvement in performance still lags
the better oxide glasses where essentially concentration independent radiative lifetimes
have been observed up to about 2 at% Er3+ [317]. It also needs to be reinforced that
sufficient gallium was present to ensure erbium ion clustering did not occur according
to prior research [120]. The concentration quenching present is therefore more intrinsic
in nature, occurring even at low erbium ion doping and high Ga:Er3+ ratios, and reasons
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for this relatively strong dependence contrary to those in some oxide glasses remain to
be elucidated.
Both the ‘zero power’ and high power 1/e lifetimes show a more significant
dependence on concentration and can be fitted well by the empirical formula mentioned
above as shown in Figure 6.8. The Q value, referred to as quenching concentration, is
1.1 for the zero pump power case and 0.545 for the high pump power, based on the
fitting. This indicates that mechanisms that depend upon the inverted population (co-
operative upconversion and/or excited state absorption process for example) are
strongly active in these glasses. The relatively strong power dependence indicates that
care is required in the design of amplifier devices using these materials where the pump
intensity has to be considered relative to the erbium ion concentration, and the likely
trade-offs thus will impose. Of particular concern is the considerable lifetime shortening
at high power and high concentrations, the region of operation often used in planar
waveguide amplifiers. Further study is needed to determine the dominant mechanism at
play and potential remedies for this behaviour. The high power lifetime shortening in
particular suggests that concentrations below 1 at% Er3+ are strongly preferable for
waveguide amplifier devices, though a trade-off with the propagation loss then will
determine the optimum concentration and design. Concentrations in the range of 0.5-
0.75 at% Er3+ might be considered optimum as they would be expected to show up to
about a factor of two reduction in the 1/e lifetime under high pumping compared to the
radiative value (which is tolerable), and this also lies in the high efficiency part of the
PL intensity versus pump power density curve as shown in Figure 6.5.
6.1.6 Erbium ion doped Ge25Ga10Se65 bulk glass summary
Broad PL emission band centred at 1538 nm was observed from a series of Er3+:
Ge25Ga10Se65 bulk glasses with erbium ion concentration ranging from 0.1 at% to 2.0
at%. The dependence of PL intensity and lifetime with increasing erbium ion
concentration was presented and discussed, clear concentration dependences being
present leading to continuous monotonic decay of the lifetime with increasing erbium
ion concentration. Based on the results, a concentration in the range of 0.5-0.75 at%
could be considered as optimum, and in the following film fabrication part, the erbium
ion concentration was controlled carefully to reside in this range.
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6.2 Erbium ion doped Ge-Ga-Se film deposition and
characterization
6.2.1 Erbium ion doped Ge-Ga-Se film deposition
Despite erbium ion doped Ge-Ga-Se glasses being intensively studied, only a few
reports on erbium ion doped Ga-containing chalcogenide films were published so far.
Takahiko et al. reported the properties of Er3+ doped Ga-Ge-Se films on fused silica
substrates deposited using RF sputtering, and lifetimes of 1.8-2.6 ms for the Er3+ 4I13/2
state were observed when excited at 973 nm [219]. In Nazabal et al.’s reports, Er3+
doped Ga-Ge-Sb-S(Se) films were fabricated using pulsed laser deposition (PLD) and
RF sputtering, and both physical and optical properties of the films were investigated
[318]. The lifetime of the 4I13/2 state of Er3+ decreased from 1.6 ms to 1.1 ms in the
sputtered films as the erbium ion concentration increased from 0.3 to 1.5 at% (3.4x1019
to 1.65x1020 ions/cm3) without mentioning the pumping wavelength. A lifetime of 1.8
ms (equal to the values calculated using Judd-Ofelt theory) was achieved on a piece of
Ga-Ge-Sb-S(Se) bulk glass with 0.05 at% Er3+ concentration after 60 minutes annealing
[318].
A large part of the difficulty in film/waveguide amplifier fabrication, particularly in
gallium containing chalcogenide glasses, relates to the difficulties in fabricating high
quality erbium ion doped films. Thermal evaporation is perhaps the simplest, and
therefore, most widely used method of preparing chalcogenide glasses films. However,
many ternary and quaternary chalcogenide glasses display the undesirable property of
forming phase separated molecular liquids on melting, with the different phases boiling
off at different temperatures and rates. This can result in films with often quite different
composition to the starting materials e.g. [213, 225]. Barely un-doped films obtained
from rare-earth ion doped bulk targets using thermal evaporation were mentioned in
Fick and Lyubin’s papers [128, 130]. The weight percentage of gallium was found to be
significantly reduced from the value in the starting material in films deposited using
standard thermal evaporation [319]. Among all the chalcogenide glass film deposition
methods, elemental co-thermal evaporation provides a solution to this issue as it offers
the possibility of controlling the evaporation rate of each element in the film
independently. Thus, the final composition of the film can be controlled precisely.
For this reason, co-thermal evaporation was employed in this work, where each
element had its own source and evaporation rate monitor, and thus the evaporation rate
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of each element could be controlled accurately and independently to select any desired
film composition. High purity Ge, Ga, Se and Er elements (5N) were used as starting
materials, and evaporation was performed at a vacuum level ~1x10-5 Pa using the
previously described chamber manufactured by Angstrom Sciences (in Section 3.2).
Procedure of the Er3+:Ge-Ga-Se film evaporation was similar to the procedure of
Er3+:As2S3 film thermal evaporation, which was described in Section 3.2.3. The film
thickness and linear refractive index (n) were measured using a dual angle spectroscopic
reflectometer (SCI FilmTek 4000) using a Tauc-Lorentz model. The final film
composition was determined by energy dispersive X-ray spectroscopy (EDS).
6.2.2 Erbium ion doped Ge-Ga-Se film characterisation
A film with a thickness of 1062±5 nm and refractive index of 2.433±0.002 at 1550 nm
was deposited. Its composition was measured to be 24.60 at% Ge, 10.94 at% Ga, 63.74
at% Se and 0.71 at% (~2×1020 ion/cm3) Er which is in the optimal region based on the
PL and lifetime measurements of the bulk glasses as noted in Section 6.1.6. An issue
encountered during film evaporations was “spitting” of gallium particles out of the
evaporation crucible. As will be discussed in this section, this led to a significant density
of small particles in the films.
The Raman spectrum of the obtained films was measured with a Horiba Jobin Yvon
64000 spectrometer utilizing a 632 nm laser as the excitation source and a 50x near
infrared objective with NA of 0.75. The obtained spectrum of an Er3+:Ge-Ga-Se film
and a reference Raman spectrum of a 0.5 at% Er3+:Ge25Ga10Se65 bulk glass are shown in
Figure 6.9. Although the spectrum from the film is much noisier than that from the bulk
glass due to the low film thickness, the main features at 202 cm-1 with a shoulder near
217 cm-1 and the lower intensity broad band from 270-330 cm-1 remain unchanged,
implying the microstructure in film is similar to its bulk counterpart. The small extended
band at ~450 cm-1 of the film arises from the thermally oxidized silicon (TOX) substrate
upon which the film was deposited.
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Figure 6.9 Raman spectra of co-thermal evaporated Er3+:Ge-Ga-Se film and 0.5 at%
Er3+:Ge25Ga10Se65 bulk glass with excitation at 632nm.
The lifetime of the 4I13/2 state of Er3+ of this Er3+:Ge-Ga-Se film was measured using
the all fibre confocal set up mentioned in Section 3.3.2 with a 1490 nm pump. The
intrinsic lifetime of this film was 0.87 ms, which is somewhat shorter than the ~1.35 ms
observed in the corresponding bulk glass. Several factors could account for this. Firstly,
the film was fabricated under non-equilibrium conditions in which a population of
homopolar and ‘wrong’ bonds could be created, and these ‘wrong’ bonds may change
the local environment for the erbium ions thereby degrading the performance. Secondly,
local erbium ion clusters may also be formed during the evaporation as pure erbium
metal was used as a source (i.e. polyatomic evaporation could occur as discussed with
Er3+:As2S3 in Section 3.5), leading to quenching of the emission [131]. Thirdly, the
nanoscale homogeneity of the films was uncertain compared to the bulk glasses. The
bulk glasses which were quenched from a metastable equilibrium melt were expected to
be located in a region of the phase diagram where the glass was essentially single phase,
rather than nanoscale phase separated. Given the clear issues with the gallium source
and the non-equilibrium nature of the film growth, the film homogeneity cannot be
guaranteed, and the “granular” nature often observed in evaporated thin films may also
be relevant here [320]. Lastly, the erbium metal oxidisation and impurities introduced in
the film during film evaporation may also cause a shorter lifetime.
The dependence of PL intensity and 1/e lifetime of Er3+ on pump power was also
investigated in the obtained Er3+:Ge-Ga-Se films, which is shown in Figure 6.10(a). The
emission has approximately quadratic pump power dependence in the measured range,
which is similar to the bulk glasses. The 1/e lifetime of the 4I13/2 state of Er3+ in Ge-Ga-
Se film also decay with increase of pump intensity as seen in the bulk glasses. To
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compare the 1/e lifetime of film and bulk glass, the observed lifetime empirical formula
(equation 6.1) was employed again to estimate the 1/e lifetime of a bulk sample with 0.7
at% Er3+ at a pump intensity around 300 kW/cm2. This produced an expected 1/e
lifetime in the 0.7 at% Er3+ doped glass of 0.74 ms compared to the 0.52 ms measured
in the film. The lifetime reduction in the film under pumping is smaller than that seen in
the radiative lifetime (27% vs. 35%), and the shape of the decay curve in Figure 6.10(b)
indicates that further reductions at the highest pump intensities expected in a typical
waveguide device (up to ~1MW/cm2) will be modest leaving a workable lifetime
around 0.5 ms.
Figure 6.10 PL intensity (a), and 1/e lifetime of the 4I13/2 state of Er3+ in the Er3+:Ge-Ga-Se
film (b) versus pump intensity in co-thermal evaporated films.
Measurements of film loss and erbium ion absorption in this Er3+:Ge-Ga-Se film
were performed using a Metricon 2010 prism coupler, and the detail of this method can
be found in Section 3.3.2. The resulting optical loss of the film, as well as a fitted
absorption curve based on the absorption curve of the Er3+:Ge25Ga10Se65 (0.1 at%) bulk
glass are shown in Figure 6.11.
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Figure 6.11 Absorption spectrum of Er3+ doped Ge-Ga-Se films measured using prism
coupler, with a fitted curve based on the absorption curve of the Er3+:Ge25Ga10Se65 (0.1
at%) bulk glass.
A minimum value of 0.8 dB/cm at 1650 nm is found outside the erbium ion
absorption band. Given the observed particles in the as deposited film, it is likely that
the loss is dominated by scattering off particulates. To investigate the film quality and
particle induced scattering loss, the deposited film was inspected using dark field
microscopy with x20 and x100 magnifications to estimate the particle density and size.
Optical micrographs of these particles taken in dark field is shown in Fig 6.12. Custom
software was used to count the particles in numerous images to ascertain an average
particle density and also to estimate the size distribution. The density of particles was
~0.01 particles/μm2, with a mean particle diameter from 1.0 to 1.1 μm. The size
distribution was always tightly bounded though with varying shape but the upper and
lower bounds were 0.9 and 1.15 µm diameter based on multiple measurements of
different parts of the wafer. Mie scattering would be expected to dominate in this size
range, and here the r/ (r is the radius of particle, and is the wavelength) ratio ranges
from 0.3 in the 1550 nm band up to around 0.5 at 1000 nm. Mie scattering is a complex
phenomenon requiring careful modelling for accurate predictions, but using the example
of water droplets in air [321], this range of size parameter would result in a scattering
cross-section that reduced by about a factor of two going from 1000 nm to 1550 nm.
Additionally, the density of particles (i.e. a 1 mm wide beam typical of the set-up used
would encounter 10,000 particles per mm of propagation length) and their likely
metallic nature of the gallium droplets would be expected to induce non-negligible loss.
It is expected that the particles could be eliminated with further refinement of the
148
evaporation set-up, most likely by using a baffled evaporation source that has no line of
sight between the evaporant and the wafer.
Figure 6.12 Particles on the Er3+ doped Ge-Ga-Se film, which was taken with microscopy
in dark field mode. The micro-sized particles has diameters ranging from 0.9 μm to 1.15
μm.
As it was established that the particles from gallium spitting cause a significant part
of the propagation loss of the film, a modification of the film evaporation facility was
implemented to resolve this issue and will be detailed in the following section.
6.3 Erbium ion doped Ge-Ga-Se film modification and
waveguide characterisation
6.3.1 Modification on erbium ion doped Ge-Ga-Se film deposition
So far, erbium ion doped Ge-Ga-Se film with controllable film composition was
achieved by co-thermal evaporation. Intrinsic lifetime of 0.9 ms of the 4I13/2 state of Er3+
in a 0.7 at% erbium ion doped Ge-Ga-Se film showed the potential erbium ion solubility
of this host material. However, a high density of particles with diameter of ~1 µm was
observed on the film surface, which resulted in non-negligible propagation loss. The
formation of these particles was suspected to be due to the spitting from gallium source.
To eliminate the particulates in the Er3+:Ge-Ga-Se film, a refinement of the
evaporation set-up was proposed. A baffled source that had no direct line of sight
between the evaporant and the wafer was used. After Er3+:Ge-Ga-Se film deposition, the
film needed to be patterned into waveguides using dry etching, the difficulty in directly
etching erbium ion containing films due to the involatile erbium compounds, as
described in Section 4.2.1, still existing. Therefore, a top layer of As2S3 was added on
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the Er3+:Ge-Ga-Se film and then etched to get the designed ridge structure see Figure
6.16(a).
Erbium ion doped Ge-Ga-Se films were deposited using co-thermal evaporation on
100 mm Thermally Oxidised Silicon wafers at room temperature in a vacuum of ~10-5
Pa. A strip loaded geometry was adopted using a 450 nm As2S3 layer atop a 650 nm
erbium ion doped Ge-Ga-Se layer. After deposition, this bilayer film was subjected to
thermal post-treatment at 130 °C for 24 hours under vacuum to bring the As2S3 film
back closer to its bulk state. Thickness and refractive index of this two layer-film were
then measured using a spectroscopic reflectometer (SCI FilmTek 4000) using a Tauc-
Lorentz model. The first 650 nm erbium ion doped Ge-Ga-Se layer had a refractive
index of 2.43 at wavelength of 1550 nm, while the 450 nm As2S3 layer had a refractive
index of 2.41 at 1550 nm after thermal treatment. This minor difference in refractive
index made it possible to design and fabricate ridge waveguide.
The obtained erbium ion doped Ge-Ga-Se film was carefully inspected under an
optical microscope using an x100 objective in dark field mode. A significant reduction
in particle density was observed, indicating the baffled furnace effectively stopped the
gallium spiting that was previously problematic in Section 6.2.2.
The composition of the Ge-Ga-Se layer was determined as Ge2.6Ga27.7Se69.7 (at%) by
EDS. This was a considerable difference to the film obtained previously and the desired
composition (Ge25Ga11Se64 at%). It was not possible to tune back to the desired
composition as the Ge furnace was operating close to its temperature limit and so the Ge
flux could not be increased, and slowing the rate of the other evaporants was not
possible as the Er rate became unstable at the low values required for the desired doping
after reducing the rates of the other furnaces. Due to the limited compositional
resolution of EDS, the low erbium ion concentration could not be read accurately from
the EDS results directly. However, comparing the optical erbium ion absorption band in
the final waveguides with an erbium ion doped films with a similar host, the erbium ion
concentration in this film was estimated at 0.5 at% (~1.5×1020 ion/cm3).
The all-fibre confocal set-up as described previously (see Section 3.3.2) was
employed to measure the PL intensity and the lifetime of the 4I13/2 excited state with
laser excitation at 1490 nm. An intrinsic lifetime of 0.99 ms was obtained from this film
under an excitation intensity of 350 kW/cm2. The dependence of lifetime and PL
intensity versus pump power had also been investigated for this film with excitation at
1490 nm, and the results are shown in Figure 6.15. Intrinsic lifetime remained at 0.99
150
ms showing no dependence with pump intensity ranging from 15 kW/cm2 to 1100
kW/cm2 as expected. The 1/e lifetime dropped from 0.88 ms at 15 kW/cm2 to 0.80 ms at
1100 kW/cm2 pump intensity, indicating interactions between pump photons and
excited ions or ion-ion effects, but the 1/e lifetime even at low pump powers, is only
slightly less than the intrinsic lifetime. Whilst a difference between the 1/e and intrinsic
lifetimes was observed in bulk glasses at this concentration, in the film sample the
difference was greater than the bulk glass. This likely indicates that either other parasitic
effects are occurring or that the erbium ions are not homogeneously distributed in the
film. Nevertheless, the lifetime was sufficient to build a reasonable amplifier. The
approximately quadratic trend of PL intensity versus pump intensity shown in Figure
6.15(b), also implies either or both of the occurrence of energy exchange effects and/or
the PL intensity saturation under high pumping intensity.
Figure 6.15 1/e and intrinsic lifetime of the 4I13/2 state of Er3+ (a); PL intensity (b) versus
pump intensity at 1490 nm of a Er3+:Ge-Ga-Se film.
A longer intrinsic lifetime of 1.34 ms could be achieved by thermally post-treating
Er3+ doped Ge-Ga-Se films at 280 °C for 24 hours, indicating the erbium ion was going
into solution better as expected from the solubility of erbium ion in Ga-containing hosts
[119, 198]. The Tg of the obtained film could not be directly measured, but from the
literature, glasses with similar composition (Ge5Ga20Se75 (at%)) had a Tg of 270 °C
[322], and Tg increased with increasing Ga content due to rise in mean bond energy of
the system. Thus, it was expected that film used in this experiment had a Tg higher than
270 °C. The large Tg difference between the layers (As2S3 layer had Tg of 180 °C),
limited annealing to only 130 °C to protect the As2S3 layer and consequently potentially
lowered Er3+ performance in the waveguides [133].
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6.3.2 Erbium ion doped Ge-Ga-Se waveguides characterization
The ridge waveguides were structured using contact lithography with AZ MIR 701
positive Photoresist and Reactive Ion Etching (RIE) with CHF3 gas [13]. 400 nm of the
450 nm As2S3 was etched away without touching the erbium ion doped layer to provide
a smooth top surface. The cross-section of the designed structure for a 2 μm waveguide
is shown in Figure 6.16. With this structure, the overlaps for both the TE and the TM
fundamental modes and the Er3+ doped area were around 70%, thereby promising
effective use of both the pump energy and the excited erbium ions.
Figure 6.16 Details of the designed structure (a), and the simulated TE fundamental mode
for the 2 μm wide Er3+:Ga-Ge-Se waveguide (b).
Cut-back measurements were performed in the TE mode at 1650 nm which was
away from the erbium ion absorption band to ensure the accuracy of the propagation
loss. The results of four waveguides from the same set with 2 µm width are shown in
Figure 6.17. The waveguide was originally 63 mm long and was then cleaved into a 22
mm section and a 41 mm section. The minimum loss among the 2 µm waveguides is
1.67 dB/cm, while the average number for these four waveguides is 1.97 dB/cm with a
standard deviation of 0.23 dB/cm. Two loss data points from 22 mm long waveguides
are absent in the figure, due to a bad cleaved edge that could be clearly seen under
microscope. It is worth noting that this propagation loss value is much higher than the
typical 0.35 dB/cm from As2S3 waveguides (in Section 4.5.1) with similar dimensions
[133].
152
Figure 6.17 Cut-back method for the loss measurement of 2 μm wide Er3+:Ge-Ga-Se
waveguides in the TE mode at 1650 nm.
The wavelength dependence of the propagation loss in a 2 µm wide Er3+:Ge-Ga-Se
waveguide was measured using a supercontinuum source and optical spectrum analyser
and is shown in Figure 6.18. As stated in Section 4.5.1, propagation loss following a
1/λ2 dependence is expected for sidewall scattering (λ is the wavelength), whilst a 1/λ4
dependence is indicative of scattering off nanoscale inhomogeneities as Rayleigh
scattering. Good propagation loss fitting was obtained for the measured waveguides
with the 1/λ2 formula, indicating the sidewall roughness in the film was the main
contributor to propagation loss. The sidewall scattering component was higher than
“normal” for one clear reason and one likely option. This waveguide had an air top
cladding, resulting in larger scattering losses from the larger refractive index difference
compared the polysiloxane clad (n=1.535 at 1550 nm) all As2S3 devices. Sidewall
scattering is often approximated as scaling with the square of the core-clad index
difference [323], which would implied a sidewall roughness loss of ~0.8 dB/cm might
be expected in this device if the etch quality was up to the standards of the all As2S3
waveguides and scaling from the 0.3 dB/cm figure there.
Nanoscale phase separation is one process that is known to occur in chalcogenide
glasses [250] that can produce the nanoscale inhomogeneities that manifest as Rayleigh
scattering. Erbium ion clustering or the growth of different structures around the erbium
ion itself can also produce similar effects as was observed in the case of previous work
on Er3+:As2S3 film (see Section 4.5.1). In waveguides measured in this work, the
Rayleigh scattering component is so small that is almost absent from our fitting results
shown in Figure 6.18, indicating there are far less film growth issues. Also as will be
shown shortly, based upon the pumping results obtained from this waveguide, little of
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the erbium ions are clustered in this film and this is consistent with the result of the
above analysis that scattering caused by doped erbium ion is unlikely in this waveguide.
Therefore it appears that the etching in this device is not as good, and in fact given the
square dependence on roughness and the acute sensitivity to the correlation length in
this region of operation [324], then only a small change is required to degrade the loss
to the levels observed.
Figure 6.18 Measured optical propagation loss spectrum of a 2 μm wide Er3+:Ge-Ga-Se
waveguide with 1/λ2 fitted data.
6.4 Optical enhancement and photoinduced absorption of
erbium ion doped Ge-Ga-Se waveguides
Excitation experiments were performed on an Er3+:Ge-Ga-Se waveguide. Given the low
Er3+ absorption at 1490 nm evident in Figure 6.18, a diode laser operating at 1505 nm
was employed as an excitation source. The scheme of the optical set-up for this
experiment is similar to the one shown in Figure 4.20. A heavily attenuated
supercontinuum source (<-30 dBm to avoid exciting the erbium ions) with range from
600-2000 nm was used as the signal source. The pump source and signal were
combined together through a 90/10 coupler (pump on 90% port) and then delivered to a
26 mm long Er3+:Ge-Ga-Se waveguide via a lensed fibre with a 2.5 µm 1/e2 diameter.
The emerging light was collected by another lensed fibre and then recorded on an
optical spectrum analyser (Ando AQ 6317). The resulting optical enhancement curve
and its dependence on excitation power are shown in Figure 6.19 as a function of pump
power.
154
Figure 6.19 Erbium ion absorption curve and optical enhancement spectra of an Er3+:Ge-
Ga-Se waveguide under different excitation intensity at 1505 nm.
From the optical enhancement curve in Figure 6.19, about 9 dB enhancement is
achieved with the available maximum 130 mW pump power at the input of the lensed
fibre (and ~50 mW launched into the waveguide). However, it is clear from looking at
the data at 1620 nm, where no enhancement would be expected, that the whole level of
the curve drops by up to 4 dB at maximum pump power. Increased loss is also seen
across the whole transmission spectrum of the device and so does not just represent
excited state absorption on the long wavelength side of the erbium ion absorption. This
loss was completely and relatively rapidly reversible (~1 minute) by turning off the
pump, and was also verified to be associated with the erbium ion as injecting high
power at 1430 nm (where the erbium ion absorption was negligible) had no effect.
The correlation with pumping the erbium ion strongly suggested an effect connected
with upconversion to shorter wavelengths, especially as such processes will be
enhanced in chalcogenide glasses compared to oxide hosts due to the much longer upper
state lifetimes resulting from much lower multiphonon recombination rates [117].
Illumination with light at and above the band gap has previously been established to
result in photoinduced wideband absorption in some materials [325]. Whilst most of the
previous study was on As2S3 and As-S-Se based fibre and showed relatively weak
absorption at 1550 nm (<0.2 dB/m at 10 mW/cm2 of transverse bandgap light [325]),
effects were also noted in Ge-As-Se glasses but little data presented. With this in mind,
1550 nm light was propagated through waveguides made from un-doped Ge3Ga28Se69,
As2S3, Ge33As12Se55 (AMTIR-1, unannealed) and Ge11.5As23.5Se65 (unannealed) whilst
exposing them to ~10 mW/cm2 red and green LED based light from above. The As2S3
155
waveguide exhibited a change of <0.01 dB/cm propagation loss, whereas all the Ge
containing waveguides showed much stronger effects. The effect was most dramatic in
the Ge33As12Se55 waveguide which exhibited an increase in propagation losses
approaching 10 dB/cm, compared to ~0.5 dB/cm in the Ge11.5As23.5Se65 and ~1 dB/cm
in the Ge3Ga28Se69 for red illumination and ~2 dB/cm for green. The bandgap of this
erbium ion doped Ge-Ga-Se film is about 1.75 eV (708 nm) based upon the results from
FilmTek measurements. The photo-induced loss change in a 28 mm long Ge-Ga-Se
waveguide with green light was recorded and is displayed in Figure 6.20. At point A,
the green light LED (~550 nm) with ~10 mW/cm2 intensity was turned on, and at point
B, the LED was turned off. The time response of the photoinduced absorption in the
materials was best fitted with a triple exponential decay. For the Ge-As-Se glasses,
recovery was completed within 2 seconds, whereas in the Ge-Ga-Se glass complete
recovery required ~90 s with fitted exponential time constants of ~0.1 s, ~1.5 s, and ~30
s. In all cases the presence of a fast decay component indicating a free carrier
component was checked for, but no response faster than ~50 ms could be found even
when looking with a 1MHz bandwidth detection system. The much increased decay
time may indicate a higher density of defect states in the sub bandgap region where the
carriers move between them as a result of thermalisation thereby surviving longer.
Figure 6.20 Photo-induced loss (at 1550 nm) change with time in a 28 mm length Ge-Ga-Se
waveguide. At point A the green light was turned on, and at point B the green light was
turned off.
Up-conversion related emissions at 980 nm, 800 nm, 670 nm, 540 nm and 520 nm,
associated with radiative decay from the 4I9/2, 4F9/2,
4S3/2 and 2H11/2 states, respectively,
to the ground state, are all located in the bandgap or Urbach tail of this Ge-Ga-Se host
material. Due to the small mode area (~2 µm2) of the waveguides, the intensity of up-
conversion related emissions in the waveguides could certainly be high enough to
156
induce photoinduced absorption, and this coupled with the insensitivity to high power
1430 nm light suggested strongly this is the cause of the increased loss. In addition, Er3+
excited to states with energy equal to or higher than the bandgap energy of the glass
matrix may directly transfer energy to the electronic states of the glass matrix [147,
148]. This may also be a contributor for the observed photoinduced loss.
To obtain an accurate erbium ion population inversion ratio, the effects of
photoinduced loss should be eliminated from the calculation of enhancement. Andriesh
[326] found an exponential fit to the photoinduced absorption in the above bandgap and
Urbach tail region, but the dependence then became non-exponential beyond ~1100 nm.
Given that the data being used was in the >1100 nm region and that a different material
was being considered, a phenomenological approach was adopted. The photoinduced
losses were extracted from the wideband difference between the zero-excitation and 130
mW-excitation curves and empirically found to be well fitted by the formula:
loss=k⁄(λ4+b), where k and b are fitting coefficients and λ is the wavelength. At this
stage no interpretation could be placed on this dependence. Figure 6.21 illustrates the
excellent fit afforded by this relationship.
Figure 6.21 Loss increment curve of an Er3+:Ge-Ga-Se waveguide due to photoinduced
absorption effect and the fitted curve.
With the fitted photoinduced loss curve in Figure 6.21, the real signal enhancement
curve under 4100 kW/cm2 pumping was obtained and is shown in Figure 6.22(a) as the
solid line. Optical enhancements corresponding to different population inversion ratios
from 10% up to 65% were calculated and is also shown in Figure 6.22(a) as dashed
lines. From these curves, the experimental enhancement curve fits with the calculated
line having 50% population inversion. This is far higher than obtained in the best
previous measurement and would produce internal gain (see in Section 4.5.3) across
most of the spectrum, as shown in Figure 6.22(b). The sharp peak and big power
157
fluctuation around 1505 nm in Figure 6.22(b) is the residual pump power centred at
1505 nm. The maximum inversion versus wavelength curve is plotted in Figure 6.23(a),
which was calculated according to the method described in Section 2.2.2. The
maximum possible inversion calculated for 1505 nm pumping is 65%. To attain
reasonable gain requires an inversion of 60-65% as is evident from Figure 6.22(a) where
more than another 4 dB of enhancement is available at the peak, and of course more
actual gain with higher doping.
Figure 6.22 Calculated optical enhancement as a function of the population inversion and
experimental optical enhancement (a); experimental optical enhancement and the erbium
ion absorption (b) of a 26 mm Er3+:Ge-Ga-Se waveguide.
Figure 6.23 Calculated maximum pump efficiency versus pump wavelength (a); measured
peak enhancement at 1538 nm as a function of pump power (measured at fibre connector)
(b) of a 26 mm Er3+:Ge-Ga-Se waveguide.
Peak enhancement at 1538 nm was extracted from each measured enhancement
curve in Figure 6.19 and its dependence on pump power is plotted in Figure 6.23 (b). It
is clear that the saturation region has almost been reached and further increases in the
enhancement will be small (~1 dB) corresponding to a further small increase in the
inversion to perhaps 55%. Note that with overlap loss and reflection the expected power
coupled into the waveguide was ~50 mW and so it was not pumped especially strongly
158
but no further pump power was available to push it to saturation. This is, to the Author’s
knowledge, the first time population inversion at these levels in erbium ion has been
attained in a chalcogenide glass host waveguide, and indicates that sufficient population
inversion for high gain amplification can be achieved with erbium ions. The saturation
at levels below the theoretical maximum indicates that a small proportion of erbium
ions are still clustered, in energy contact, or are optically inactive in this host material as
grown, consistent with the low power lifetime shortening compared to the intrinsic
lifetime noted in Section 6.3.1. Further improvement in both deposition and/or more
effective thermal post-treatment is required to match the theoretical performance. This,
however, requires a change of the waveguide design to use a higher Tg strip glass, or the
use of waveguide fabrication methods such as lift off [274] or hot embossing [280, 327].
A practical device also requires much lower passive waveguide losses (should be
achievable with improved processing and film growth with appropriate compositions)
and also the resolution of the photoinduced losses. It remains to be seen whether the
photoinduced loss also occurs in films with the desired composition of gallium doped
germanium selenide, or whether a different composition with a higher bandgap will be
more advantageous.
6.5 Conclusion
In this chapter, high quality erbium ion doped Ge-Ga-Se films were deposited and ridge
waveguides based on these films were patterned. Significant signal enhancement at 1.5
µm was observed and 50% erbium ion population inversion was obtained, with a
saturation maximum of ~55% being possible, in waveguides with 1.5×1020 Er3+
ion/cm3. This is the highest level of inversion ever demonstrated for erbium ions in a
chalcogenide glass host waveguide and is an important step towards future devices
operating at 1550 nm and on the MIR transitions in erbium ion. Photoinduced
absorption loss caused by upconversion products in the waveguides was the remaining
hurdle to achieving net gain. Further research is needed to find suitable compositions
that possess high rare-earth ion solubility whilst avoiding the detrimental photoinduced
losses.
159
Chapter 7
Conclusions and recommendations
7.1 Conclusions
This work focused on the fundamentals of fabrication and characterisation of erbium ion
doped chalcogenide glass planar waveguide amplifiers. There have been a number of
important achievements that advance the field. Erbium ion doped chalcogenide films
were deposited using co-thermal evaporation and RF sputtering based on As2S3 and Ge-
Ga-Se host materials. Waveguides were patterned with standard photolithography and
reactive ion plasma etching. In erbium ion doped As2S3 waveguides made using co-
thermal evaporation, internal gain from 1570 to 1630 nm was observed for the first time
in a chalcogenide glass host. Erbium ion doped As2S3 films deposited through RF
sputtering showed even more promising results in terms of PL lifetime and PL intensity
of Er3+. Significant signal enhancement at 1.5 µm was observed and 50% erbium ion
population inversion was obtained in an erbium ion doped Ge-Ga-Se waveguides
fabricated using co-thermal evaporated film. To the Author’s knowledge this is the
highest level of inversion ever demonstrated for erbium ions in a chalcogenide host and
is an important step towards future devices operating at 1.5 µm and on the MIR
transitions in erbium ion in chalcogenide planar devices.
This work had extensively investigated the fabrication methods for high quality
erbium ion doped chalcogenide thin films. Co-thermal evaporation was employed in
erbium ion doped As2S3 deposition for its precise control of each element’s evaporation
rate and thus the final film composition. Both physical and optical properties of
obtained film were investigated carefully. Films with erbium ion concentration from
0.15 at% (0.45x1020 ions/cm3) to 0.6 at% (1.8x1020 ions/cm3) were deposited and
studied to optimise the erbium ion concentration. Film doped with 0.15 at% Er3+ was
shown to have the best performance in terms of the 4I13/2 state lifetime and 1.5 µm PL
efficiency. However, results from high temperature thermal post-treatment showed the
solubility of erbium ion in As2S3 film is even lower than 0.15 at%. Likewise it was
shown that the clustering of erbium ion in the films resulted from polyatomic
evaporation, a difficult matter to resolve using liquid metal sources whilst still making
highly doped films in an acceptable time frame. With the low loss waveguide (~0.35
160
dB/cm) patterned on 0.15 at% Er3+ doped As2S3 film, internal gain was observed from
1570 to 1630 nm for the first time.
Different approaches to fabricate erbium ion doped As2S3 films were investigated.
Promising results on intrinsic lifetime were achieved in ErCl3 doped As2S3 (2.6 ms),
however no waveguide was fabricated on these films due to poor film quality. Films
fabricated using RF sputtering showed promising lifetime (2 ms after thermal post-
treatment) and strong PL intensity (stronger than the PL intensity from the previous co-
thermal Er3+:As2S3 film with internal gain at similar erbium ion concentration), but its
emission performance was sensitive to green light and moisture. Due to the existence of
green light from up-conversion in erbium ion doped waveguide when high pumping
intensity was applied, the PL performance of this sputtered film was unstable and
decayed to a level that is not suitable for an amplification device.
Host chalcogenide material other than As2S3, Ge-Ga-Se glass, was also investigated
for its high rare-earth ion solubility property. Emission properties of Er3+ doped Ge-Ga-
Se bulk glasses were studied as a function of erbium ion doping concentration from 0.1
at% up to 2 at%. A clear concentration quenching effect in the Ge-Ga-Se glasses was
observed for the first time even when sufficient gallium was present to ensure
homogeneous distribution of the erbium ions. A region between approximately 0.5 and
0.8 at% of Er3+, however, was shown to provide sufficient doping, good
photoluminescence and adequate lifetime to envisage practical planar waveguide
amplifier devices. Micron-size particles in erbium ion doped Ge-Ga-Se caused by
spitting of the gallium was observed and eliminated using a modification to the
deposition equipment. High quality erbium ion doped Ge-Ga-Se films were achieved
and ridge waveguides based on these films were patterned. Significant signal
enhancement at 1.5 µm was observed and ~50% Er3+ population inversion was obtained,
with a saturation maximum of ~55% being possible in waveguides with 1.5×1020 Er3+
ion/cm3. To the Author’s knowledge this is the highest level of population inversion
ever demonstrated for erbium ion in a chalcogenide glass host. However, photoinduced
absorption loss caused by upconversion products in the waveguides showed up and
remained a hurdle to achieving net gain.
7.2 Recommendations for future work
161
Despite achieving internal gain in erbium ion doped As2S3 by co-thermal evaporation,
and more than 50% population inversion in erbium ion doped Ge-Ga-Se host, so far no
practical planar waveguide amplifier was realised. On the positive side, according to the
results obtained, it is reasonable to conclude that there is no intrinsic barrier in making
rare-earth ion doped chalcogenide waveguide amplifiers. The problem to be resolved is
in attaining unclustered erbium ion doping and moisture immunity.
Likely the best way forward for evaporated As2S3 is to move to Er2S3 as the dopant,
which requires some modification to the evaporation system so that the Er2S3 can be
evaporated from a tungsten or molybdenum boat without heating the film during
deposition. This should then conceptually prevent the erbium ion clustering during the
deposition, and if the thermal post-treatment is restricted, allow for high concentrations
to be introduced into the films and for net gain to be achieved.
Since the RF sputtered film had almost the best optical performance, excepting the
PL decay, therefore it is worth looking for a way to overcome this issue. Columnar
structure in films due to the growth habit of RF sputtering and was reported in different
host materials. Optimising the film deposition parameters such RF power and gas
pressure might be a possible way to eliminate columnar structure. A successful example
was reported in [2] in TeO2 films (though with reactive sputtering) where net gain
optical amplification was achieved. However, the As2S3 target is quite fragile, thus
cracking may occur when the RF power is high, which leads to a poor film quality.
Considering this, little space is left for increasing the RF power, but control of the gas
pressure may be a viable alternative and should be investigated. A moisture barrier
could also be employed, though work is required to establish exactly how to totally
prevent moisture penetration.
More than 50% population inversion was achieved in an erbium ion doped Ge-Ga-
Se waveguide, but strong photo-induced loss prevented internal gain being observed. It
was also noticed that with different glass compositions, the amplitude of photo-induced
loss varies considerably. For example, Ge33As12Se55 (AMTIR-1) had a much bigger
photo-induced loss than the Ge-Ga-Se glass used here, and Ge11As24Se65 (at%) had a
smaller one under the same condition. Therefore, it is reasonable to ask if there is a
possible composition in Ge-Ga-Se family that has good rare-earth ion solubility but
without significant photo-induced loss. In this case, materials research into Ge-Ga-Se to
find a better composition that has a better photo-induced loss resistance is necessary. On
the other hand, the photo-induced loss reported in Chapter 6 was thought bandgap
162
related, and the up-conversion emissions from erbium ion had a huge overlap with Ge-
Ga-Se’s bandgap which brought the issue confronted here. Thus, rare-earth ions without
any strong visible emissions and up-conversion related emissions in Ge-Ga-Se’s
bandgap wavelength may work fine with this host.
Also, the glasses this thesis focused on are only two members of the big
chalcogenide glass family. Other materials such as Ge-Ga-S, Ga-As-S, Ga-La-S etc.,
may offer better opportunity than these two, in this perspective, a comprehensive study
of chalcogenide materials is necessary.
163
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